Active electrode material

ABSTRACT

The invention relates to active electrode materials and to methods for the manufacture of active electrode materials. Such materials are of interest as active electrode materials in lithium-ion or sodium-ion batteries. The invention provides an active electrode material comprising a mixture of (a) at least one lithium titanium oxide and (b) at least one mixed niobium oxide, wherein the mixed niobium oxide is expressed by the general formula [M1] x [M2] (1-x) [Nb] y [O] z .

FIELD OF THE INVENTION

The present invention relates to active electrode materials and to methods for their manufacture. Such materials are of interest as active electrode materials in lithium-ion or sodium-ion batteries, for example as anode materials for lithium-ion batteries.

BACKGROUND

Lithium-ion (Li-ion) batteries are a commonly used type of rechargeable battery with a global market predicted to grow to $200 bn by 2030. Li-ion batteries are the technology of choice for electric vehicles that have multiple demands across technical performance to environmental impact, providing a viable pathway for a green automotive industry.

A typical lithium-ion battery is composed of multiple cells connected in series or in parallel. Each individual cell is usually composed of an anode (negative polarity electrode) and a cathode (positive polarity electrode), separated by a porous, electrically insulating membrane (called a separator), immersed into a liquid (called an electrolyte) enabling lithium ions transport.

In most systems, the electrodes are composed of an electrochemically active material—meaning that it is able to chemically react with lithium ions to store and release them reversibly in a controlled manner—mixed if necessary with an electrically conductive additive (such as carbon) and a polymeric binder. A slurry of these components is coated as a thin film on a current collector (typically a thin foil of copper or aluminium), thus forming the electrode upon drying.

In the known Li-ion battery technology, the safety limitations of graphite anodes upon battery charging is a serious impediment to its application in high-power electronics, automotive and industry. Among a wide range of potential alternatives proposed recently, lithium titanate (LTO, particularly spinel-type Li₄Ti₅O₁₂) and mixed niobium oxide-based materials are the main contenders to replace graphite as the active material of choice for high power applications.

Batteries relying on a graphitic anode are fundamentally limited in terms of charging rate. Under nominal conditions, lithium ions are inserted into the anode active material upon charging. When charging rate increases, typical graphite voltage profiles are such that there is a high risk that overpotentials lead to the potential of sites on the anode to become <0 V vs. Li/Li+, which leads to a phenomenon called lithium dendrite electroplating, whereby lithium ions instead deposit at the surface of the graphite electrode as lithium metal. This leads to irreversible loss of active lithium and hence rapid capacity fade of the cell. In some cases, these dendritic deposits can grow to such large sizes that they pierce the battery separator and lead to a short-circuit of the cell. This can trigger a catastrophic failure of the cell leading to a fire or an explosion. Accordingly, the fastest-charging batteries having graphitic anodes are limited to charging rates of 5-7 C, but often much less.

Lithium titanate (LTO) anodes do not suffer from dendrite electroplating at high charging rate thanks to their high potential (1.55 V vs. Li/Li+), and have excellent cycle life as they do not suffer from significant volume expansion of the active material upon intercalation of Li ions due to their accommodating 3D crystal structure. LTO cells are typically regarded as high safety cells for these two reasons. However, LTO is a relatively poor electronic and ionic conductor, which leads to limited capacity retention at high rate and resultant power performance, unless the material's primary particles are nanosized to increase specific surface area, and carbon-coated to increase electronic conductivity. This level of material engineering increases the porosity and specific surface area of the active material, and results in a significantly lower achievable packing density in an electrode. This is significant because it leads to low density electrodes and a higher fraction of electrochemically inactive material (e.g. binder, carbon additive), resulting in much lower gravimetric and volumetric energy densities. As such, methods that can improve the packing density such as physical mixtures of different active materials and/or particle sizes, are very attractive to improve performance.

A key measure of anode performance is the electrode volumetric capacity (mAh/cm³), that is, the amount of electric charges (that is lithium ions) that can be stored per unit volume of the anode. This is an important factor to determine the overall battery energy density on a volumetric basis (Wh/L) when combined with the cathode and appropriate cell design parameters. Electrode volumetric capacity can be approximated as the product of electrode density (g/cm³), active material specific capacity (mAh/g), and fraction of active material in the electrode. LTO anodes typically have relatively low specific capacities (c. 165 mAh/g, to be compared with c. 330 mAh/g for graphite) which, combined with their low electrode densities (typically <2.0 g/cm³) and low active material fractions (<90%) discussed above, lead to very low volumetric capacities (<300 mAh/cm³) and therefore low battery energy density and high $/kWh cost in various applications. As a result, LTO batteries/cells are generally limited to specific niche applications, despite their long cycle life, fast-charging capability, and high safety.

Mixed niobium oxides (MNO) were first identified as potential battery materials in the academic literature in the 1980's,^([2,3]) but have only seen a commercial focus since the 2010's with the demonstration of a practical cell combining a TiNb₂O₇ and a commercially-available LNMO (lithium nickel manganese oxide) cathode showing promising performance in terms of rate capability, cycle life, and energy density.^([1]) Selected MNO anodes offer characteristics that are similar to LTO in terms of high operating potential vs. Li/Li+ (1.55 V) and low volume expansion (<5%) leading to safe fast-charge and long cycle life (>10,000 cycles). A key advantage of MNO anodes is that practical specific capacities significantly higher than LTO (c. 165 mAh/g) can be achieved (c. 200-300 mAh/g), which improves cell energy density. In contrast to LTO materials (10⁻¹⁷ cm² s⁻¹), the Li-ion diffusion coefficient is typically much higher for specific MNO compositions that result in so-called “Wadsley-Roth” or “Tetragonal Tungsten Bronze” crystal structures (10⁻¹⁴-10⁻¹⁰ cm² s⁻¹).^([4]) This means that Li ions will diffuse across much greater distances through the active material within the same time for MNO materials vs LTO, at a fixed charge/discharge rate. Therefore, MNO materials can be less porous and use larger primary particles/crystals (0.5-10 μm for MNO vs <100 nm for LTO), retaining or improving the high-power charge/discharge performance. This results in higher electrode densities, and volumetric energy densities of cells, leading to a lower $/kWh cost at the application level.

However, the nominal voltage of MNO materials is typically higher than that of LTO (i.e. >1.55 V vs Li/Li⁺), which acts as a trade-off and decreases achievable energy density in full Li-ion cells. Cost of precursors, in particular Nb-based raw materials, also limits the deployment of MNO materials in commercial products for mass market applications.

US2019/0288283A1 discloses a lithium niobium composite oxide where as an essential feature some of the niobium must be replaced by at least one element selected from Fe, Mg, Al, Cu, Mn, Co, Ni, Zn, Sn, Ti, Ta, V, and Mo. The document refers to but does not exemplify an electrode comprising the lithium niobium composite oxide and another active material which may be any of lithium titanate having a ramsdellite structure, lithium titanate having a spinel structure, monoclinic titanium dioxide, anatase type titanium dioxide, rutile type titanium dioxide, a hollandite type titanium composite oxide, an orthorhombic titanium-containing composite oxide, and a monoclinic niobium titanium composite oxide.

US2020/0140339A1, US2018/0083283A1, and U.S. Pat. No. 10,096,826B2 disclose titanium niobate materials (based on TiNb₂O₇ or Ti₂Nb₁₀O₂₉). They refer to but do not exemplify mixtures with other active materials such as different forms of titanium dioxide and lithium titanate. Titanium niobate materials exhibit a typically lower lithium ion diffusion coefficient than other MNO materials.

WO2019234248A1 discloses examples of an electrode comprising a mixture of a niobium tungsten oxide (W₅Nb₁₆O₅₅) and LTO (Li₄Ti₅O₁₂). It is believed that the properties of this electrode, in particular the properties of W₅Nb₁₆O₅₅, can be improved.

The present invention has been devised in light of the above considerations.

SUMMARY OF THE INVENTION

In a first aspect, the invention provides an active electrode material comprising a mixture of (a) at least one lithium titanium oxide and (b) at least one mixed niobium oxide, wherein the mixed niobium oxide is expressed by the formula [M1]_(x)[M2]_((1-x))[Nb]_(y)[O]_(z), wherein:

-   -   M1 and M2 are different;     -   M1 is selected from one or more of P, B, Ti, Mg, V, Cr, W, Zr,         Mo, Cu, Fe, Ga, Ge, Ca, K, Ni, Co, Al, Sn, Mn, Ce, Se, Si, Sb,         Y, La, Hf, Ta, Zn, In, and Cd;     -   M2 is selected from one or more of P, Mg, V, Cr, W, Zr, Mo, Cu,         Fe, Ga, Ge, Ca, K, Ni, Co, Al, Sn, Mn, Ce, Sb, Bi, Sr, Y, La,         Hf, Zn, Ta, In, and Cd; and wherein     -   x satisfies 0≤x<0.5;     -   y satisfies 0.5≤y≤49;     -   z satisfies 4≤z≤124;     -   with the proviso that if x=0 and M2 consists of a single element         then the mixed niobium oxide is oxygen deficient.

The active electrode material is a mixture of a lithium titanium oxide and a mixed niobium oxide. The mixed niobium oxide has been modified by cation substitution and/or by introducing oxygen deficiency. The inventors have found that mixed niobium oxides that have been modified in this way have improved properties for use as active electrode materials, e.g. in anodes for lithium- and sodium-ion batteries. In particular, the inventors have found that the modified mixed niobium oxides have improved electronic conductivity, improved initial coulombic efficiency, and/or improved capacity retention at high charge/discharge rates, compared to the unmodified ‘base’ mixed niobium oxides. Therefore, the invention represents a way of combining the improved properties of modified mixed niobium oxides with the benefits of lithium titanium oxides for use as active electrode materials. In particular, the combination of mixed niobium oxides and lithium titanium oxides provide advantages versus the individual active materials with regards to improved cost, electrode formulation and ink processing, and various aspects of electrochemical performance.

The active electrode materials of the invention are particularly useful in electrodes, preferably for use in anodes for lithium-ion or sodium-ion batteries. Therefore, a further implementation of the invention is a composition comprising the active electrode material of the first aspect and at least one other component; optionally wherein the at least one other component is selected from a binder, a solvent, a conductive additive, an additional active electrode material, and mixtures thereof. Such a composition is useful for fabricating an electrode. A further implementation of the invention is an electrode comprising the active electrode material of the first aspect in electrical contact with a current collector. A further implementation of the invention is an electrochemical device comprising an anode, a cathode, and an electrolyte disposed between the anode and the cathode, wherein the anode comprises an active electrode material according to the first aspect; optionally wherein the electrochemical device is a lithium-ion battery or a sodium-ion battery.

In a second aspect, the invention provides method for making an active electrode material, wherein the active electrode material is as defined in the first aspect, the method comprising mixing at least one lithium titanium oxide with at least one mixed niobium oxide.

The invention includes the combination of the aspects and features described herein except where such a combination is clearly impermissible or expressly avoided.

SUMMARY OF THE FIGURES

The principles of the invention will now be discussed with reference to the accompanying figures in which:

FIG. 1 : XRD diffraction patterns of samples 1, 4, 14, 2, 5, 15, 16, 18 and 22;

FIG. 2 : XRD diffraction patterns of samples 8 and 9;

FIG. 3 : XRD diffraction patterns of samples 6, 7, 17, 19 and 20;

FIG. 4 : XRD diffraction patterns of samples 10, 11 and 21;

FIG. 5 : XRD diffraction patterns of samples 12 and 13;

FIG. 6 : TGA characterisation in air of sample 3;

FIG. 7 : the particle size distribution of samples 1, 2, 15, and 16;

FIG. 8 : the particle size distribution of sample 3;

FIG. 9 : SEM image of sample 3 before pyrolysis and coated with conductive Au for imaging;

FIG. 10 : SEM image of sample 3 after pyrolysis (no conductive coating);

FIG. 11 : SEM images of samples 1 and 2;

FIG. 12 : representative lithiation and delithiation voltage profiles obtained by galvanostatic cycling in half cell configuration, 1.1-3.0 V voltage window, first 2 cycles at 0.05 C rate for samples 1 and 16;

FIG. 13 : representative lithiation and delithiation voltage profiles obtained by galvanostatic cycling in half cell configuration, 1.1-3.0 V voltage window, first 2 cycles at 0.05 C rate for samples 6 and 7;

FIGS. 14 : (a) and (b) EIS measurements of samples 1, 7, and 16 at different axes scales.

FIG. 15 : particle size distributions of sample 16 before and after post-processing;

FIG. 16 : SEM image of the surface of an electrode made from sample 22, focused on the surface of an active material particle;

FIG. 17 : representative lithiation and delithiation voltage profiles obtained by galvanostatic cycling in half cell configuration, 1.1-3.0 V voltage window, first 2 cycles at 0.05 C rate for samples 12 and 13.

FIG. 18 : Powder X-ray Diffraction (XRD) spectra for Sample E1, E2, E3, and E4.

FIG. 19 : Galvanostatic charge/discharge curves at C/10, in half-cells for tests A, E and G, for the 2^(nd) lithiation/de-delithiation cycle, between 1.1-3.0 V.

FIG. 20 : Galvanostatic state-of-charge (SOC) retention at 10 C in de-lithiation in half-cells for tests B and F, between 1.1-3.0 V.

FIG. 21 : Galvanostatic de-lithiation curve at 5 C, in half-cells, for tests B and H between 1.1-3.0 V.

FIG. 22 : Galvanostatic state-of-charge (SOC) retention at 10 C in de-lithiation in half-cells for tests D and I, between 1.1-3.0 V.

FIG. 23 : Galvanostatic charge/discharge curves at C/10, in half-cells fortests A and J, for the 2^(nd) lithiation/de-delithiation cycle, between 1.1-3.0 V.

FIG. 24 : Powder X-ray Diffraction (XRD) spectra for Samples E5 and E6.

FIG. R1 shows XRD diffraction patterns of samples R1, R2.

FIG. R2 shows XRD diffraction patterns of samples R3, R4, R5.

FIG. R3 shows XRD diffraction patterns of samples R6, R7, R8.

FIG. R4 shows XRD diffraction patterns of samples R9, R10.

FIG. R5 shows the particle size distributions of samples R2, R4, R7, R10.

FIG. R6 shows representative lithiation and delithiation voltage profiles obtained by galvanostatic cycling in half cell configuration, 1.1-3.0 V voltage window, first 2 cycles at 0.1 C rate for samples R1 and R2.

FIG. R7 shows representative lithiation and delithiation voltage profiles obtained by galvanostatic cycling in half cell configuration, 1.1-3.0 V voltage window, first 2 cycles at 0.1 C rate for samples R3 and R5.

FIG. R8 shows representative lithiation and delithiation voltage profiles obtained by galvanostatic cycling in half cell configuration, 1.1-3.0 V voltage window, first 2 cycles at 0.1 C rate for samples R6 and R7.

FIG. R9 shows representative lithiation and delithiation voltage profiles obtained by galvanostatic cycling in half cell configuration, 1.1-3.0 V voltage window, first 2 cycles at 0.1 C rate for samples R9 and R10. The x axis is in terms of state-of-charge (SOC), to be able to normalise the curves to their maximum capacities and evaluate the curve shape.

FIG. R10 shows XRD diffraction patterns of R11-R14.

DETAILED DESCRIPTION OF THE INVENTION

Aspects and embodiments of the present invention will now be discussed with reference to the accompanying figures. Further aspects and embodiments will be apparent to those skilled in the art. All documents mentioned in this text are incorporated herein by reference.

The ratio by mass of (a):(b) may be in the range of 0.5:99.5 to 99.5:0.5, preferably in the range of 2 98 to 98:2. In one implementation the active electrode material comprises a higher proportion of the lithium titanium oxide than the mixed niobium oxide, e.g. the ratio by mass of (a):(b) is at least 2:1, at least 5:1, or at least 8:1. Advantageously, this allows the mixed niobium oxide to be incrementally introduced into existing electrodes based on lithium titanium oxides without requiring a large change in manufacturing techniques, providing an efficient way of improving the properties of existing electrodes. In another implementation the active electrode material comprises a higher proportion of the mixed niobium oxide than the lithium titanium oxide, e.g. such that the ratio by mass of (b):(a) is at least 2:1, at least 5:1, or at least 8:1. Advantageously, this allows for the cost of the material to be reduced by replacing some of the mixed niobium oxide with lithium titanium oxide.

Optionally, the active electrode material may consist of a mixture of (a) at least one lithium titanium oxide and (b) at least one mixed niobium oxide. Additionally, the active electrode material may consist of a mixture of (a) one lithium titanium oxide and (b) one mixed niobium oxide.

The term “mixed niobium oxide” (MNO) refers to an oxide comprising niobium and at least one other cation. MNO materials have a high redox voltage vs. Lithium (Li/Li⁺) >0.8V, enabling safe and long lifetime operation, crucial for fast charging battery cells. Moreover, niobium cations can have two redox reactions per atom, resulting in higher theoretical capacities than, for example, LTO.

The MNO is expressed by the formula [M1]_(x)[M2]_((1-x))[Nb]_(y)[O]_(z), wherein:

-   -   M1 and M2 are different;     -   M1 is selected from one or more of P, B, Ti, Mg, V, Cr, W, Zr,         Mo, Cu, Fe, Ga, Ge, Ca, K, Ni, Co, Al, Sn, Mn, Ce, Se, Si, Sb,         Y, La, Hf, Ta, Zn, In, and Cd;     -   M2 is selected from one or more of P, Mg, V, Cr, W, Zr, Mo, Cu,         Fe, Ga, Ge, Ca, K, Ni, Co, Al, Sn, Mn, Ce, Sb, Bi, Sr, Y, La,         Hf, Zn, Ta, In, and Cd; and wherein     -   x satisfies 0≤x<0.5;     -   y satisfies 0.5≤y≤49;     -   z satisfies 4≤z≤124;     -   with the proviso that if x=0 and M2 consists of a single element         then the mixed niobium oxide is oxygen deficient.

By ‘one or more of’, it is intended that either M1 or M2 may each represent two or more elements from their respective lists. An example of such a material is Ti_(0.05)W_(0.25)Mo_(0.70)Nb₁₂O₃₃. Here, M1 represents Ti_(x′)W_(x″) (where x′+x″=x), M2 represents Mo, x=0.3, y=12, z=33. Another example of such a material is Ti_(0.05)Zr_(0.05)W_(0.25)Mo_(0.65)Nb₁₂O₃₃. Here, M1 represents Ti_(x′)Zr_(x″)W_(x″) (where x′+x″+x′″=x), M2 represents Mo, x=0.35, y=12, z=33.

M2 may be selected from one or more of Mo, W, V, Zr, P, Al, Zn, Ga, Ge, Ta, Cr, Cu, K, Mg, Ni, and Hf; or one or more of Mo, W, V, Zr, P, Al, Zn, Ga, and Ge; or one or more of Mo, W, V, and Zr. Preferably, M2 consists of a single element. M2 does not represent Ti. In other words, preferably, Ti is not the major non-Nb cation in the mixed niobium oxide. Where M1 represents Ti alone, preferably x is 0.05 or less. Where M1 represents one or more cations including Ti, preferably the amount of Ti relative to the total amount of non-Nb cations is 0.05:1 or less.

M1 is a cation which substitutes for M2 in the crystal structure. M1 may be selected from one or more of P, B, Ti, Mg, V, Cr, W, Zr, Mo, Cu, Fe, Ga, Ge, K, Ni, Co, Al, Hf, Ta, and Zn; or one or more of P, B, Ti, Mg, V, Cr, W, Zr, Mo, Ga, Ge, Al, and Zn; or one or more of Ti, Zr, V, W, and Mo. M1 may have an equal or lower oxidation state than M2. Preferably, M1 has a lower oxidation state than M2. When more than one element is present as M1 and/or M2 it will be understood that the oxidation state refers to M1 and/or M2 as a whole. For example, if 25 at % of M1 is Ti and 75 at % of M1 is W the oxidation state of M1 is 0.25×4 (the contribution from Ti)+0.75×6 (the contribution from W). Advantageously, when M1 has a lower oxidation state than M2 this is compensated for by the formation of oxygen vacancies, i.e. forming an oxygen deficient mixed niobium oxide. The presence of oxygen vacancies is believed to improve the conductivity of the mixed niobium oxide and to provide further benefits, as evidenced by the examples. Optionally, M1 comprises at least one cation with a 4+ oxidation state and M2 comprises at least one cation with a 6+ oxidation state. Optionally, M1 has an oxidation state of 4+ and M2 has an oxidation state of 6+. M1 preferably has a different ionic radius than M2, most preferably a larger ionic radius. This gives rise to changing unit cell size and local distortions in crystal structure. This is believed to improve electrochemical properties such as specific capacity and Coulombic efficiency through altering the Li ion site availability by varying cavity size and reduction of energy barriers to reversible lithiation.

x defines the amount of M1 which replaces M2 in the mixed niobium oxide. Since x is <0.5, M2 is the major non-Nb cation in the mixed niobium oxide. Preferably, x>0. x may satisfy 0<x<0.5, 0.01<x<0.4, or 0.05≤x≤0.25.

y represents the amount of niobium in the mixed niobium oxide. z represents the amount of oxygen in the mixed niobium oxide. The precise values of y and z within the ranges defined may be selected to provide a charge-balanced structure. The precise values of y and z within the ranges defined may be selected to provide a charge balanced, or substantially charge balanced, crystal structure. Additionally or alternatively, the precise values of y and z within the ranges defined may be selected to provide a thermodynamically stable, or thermodynamically metastable, crystal structure, e.g. based on the unmodified crystal structures disclosed herein.

In some cases, z may be defined in the format z=(z′−z′α), where α is a non-integer value less than 1, for example where α satisfies 0≤α≤0.05. α may be greater than 0, i.e. a may satisfy 0<α≤0.05. When α is greater than 0, the mixed niobium oxide is oxygen deficient, i.e. the material has oxygen vacancies. Such a material would not have precise charge balance, but is considered to be “substantially charge balanced” as indicated above. Alternatively, α may equal 0, in which case the mixed niobium oxide is not oxygen deficient. Preferably, the mixed niobium oxide is oxygen deficient. In particular, when x=0 preferably the material is oxygen deficient.

When α is 0.05, the number of oxygen vacancies is equivalent to 5% of the total oxygen in the crystal structure. In some embodiments, a may be greater than 0.001 (0.1% oxygen vacancies), greater than 0.002 (0.2% oxygen vacancies), greater than 0.005 (0.5% oxygen vacancies), or greater than 0.01 (1% oxygen vacancies). In some embodiments, a may be less than 0.04 (4% oxygen vacancies), less than 0.03 (3% oxygen vacancies), less than 0.02 (2% oxygen vacancies), or less than 0.1 (1% oxygen vacancies). For example, a may satisfy 0.001≤α≤0.05. When the material is oxygen deficient, the electrochemical properties of the material may be improved, for example, resistance measurements may show improved conductivity in comparison to equivalent non-oxygen deficient materials. As will be understood, the percentage values expressed here are in atomic percent.

Oxygen vacancies may be formed in a mixed niobium oxide by the sub-valent substitution of a base material. For example, oxygen vacancies may be formed by substituting some of the Mo(6+) cations in MoNb₁₂O₃₃ with cations of a lower oxidation state, such as Ti(4+) and/or Zr(4+) cations. A specific example of this is the compound Ti_(0.05)Zr_(0.05)W_(0.25)Mo_(0.65)Nb₁₂O_(33-δ) which is derived from the base material MoNb₁₂O₃₃ and includes oxygen vacancies. Oxygen vacancies may also be formed by heating a mixed niobium oxide under reducing conditions (for instance, heating under nitrogen atmosphere at e.g. 800-1350° C.). A specific example of this is the compound MoNb₁₂O_(33-δ). The mixed niobium oxide may have induced oxygen deficiency. Induced oxygen deficiency may be understood to mean that the mixed niobium oxide contains additional oxygen vacancies, e.g. in addition to oxygen vacancies already present in the mixed niobium oxide due to sub-valent substitution of M2 with M1.

A number of methods exist for determining whether oxygen vacancies are present in a material. For example, Thermogravimetric Analysis (TGA) may be performed to measure the mass change of a material when heated in air atmosphere. A material comprising oxygen vacancies can increase in mass when heated in air due to the material “re-oxidising” and the oxygen vacancies being filled by oxide anions. The magnitude of the mass increase may be used to quantify the concentration of oxygen vacancies in the material, on the assumption that the mass increase occurs entirely due to the oxygen vacancies being filled. It should be noted that a material comprising oxygen vacancies may show an initial mass increase as the oxygen vacancies are filled, followed by a mass decrease at higher temperatures if the material undergoes thermal decomposition. Moreover, there may be overlapping mass loss and mass gain processes, meaning that some materials comprising oxygen vacancies may not show a mass gain (and sometimes not a mass loss or gain) during TGA analysis.

Other methods of determining whether oxygen vacancies are present include electron paramagnetic resonance (EPR), X-ray photoelectron spectroscopy (XPS, e.g. of oxygen 1s and/or and of cations in a mixed oxide), X-ray absorption near-edge structure (XANES, e.g. of cations in a mixed metal oxide), and TEM (e.g. scanning TEM (STEM) equipped with high-angle annular darkfield (HAADF) and annular bright-field (ABF) detectors). The presence of oxygen vacancies can be qualitatively determined by assessing the colour of a material relative to a non-oxygen-deficient sample of the same material. For example, stoichiometric MoNb₁₂O₃₃ has a white, off-white, or yellow colour whereas oxygen-deficient MoNb₁₂O_(33-δ) has a purple colour. The presence of vacancies can also be inferred from the properties, e.g. electrical conductivity, of a stoichiometric material compared to those of an oxygen-deficient material.

When the mixed niobium oxide is oxygen deficient it may be selected from MoNb₁₂O_((33-33α)), WNb₁₂O_((33-33α)), Mo₃Nb₁₄O_((44-44α)), VNb₉O_((25-25α)), ZrNb₂₄O_((62-62α)), Zn₂Nb₃₄O_((87-87α)), Cu₂Nb₃₄O_((87-87α)), WNb₄O_((31-31α)), W₉Nb₈O_((47-47α)), W₅Nb₁₆O_((55-55α)), W₁₆Nb₁₈O_((93-93α)), AlNb₁₁O_((29-29α)), GaNb₁₁O_((29-29α)), FeNb₁₁O_((29-29α)), AlNb₄₉O_((124-124α)), GaNb₄₉O_((124-124α)), FeNb₄₉O_((124-124α)), and GeNb₁₈O_((47-47α)) wherein α satisfies 0<α≤0.05. These are examples of materials where x=0 and M2 consists of a single element. Preferably when the mixed niobium oxide is oxygen deficient it is selected from MoNb₁₂O_((33-33α)), WNb₁₂O_((33-33α)), VNb₉O_((25-25α)), ZrNb₂₄O_((62-62α)), W₅Nb₁₆O_((55-55α)), W₇Nb₄O_((31-31α)), and W₉Nb₈O_((47-47α)) wherein α satisfies 0<α≤0.05.

The mixed niobium oxide may be selected from M1_(x)Mo_((1-x))Nb₁₂O_((33-33α)), M1_(x)W_((1-x))Nb₁₂O_((33-33α)), M1_(x)Mo_((1-x))Nb_(4.667)O_((14.567-14.667α)) (i.e. Mo₃Nb₁₄O₄₄ base structure), M1_(x)V_((1-x))Nb₉O_((25-25α)), M1_(x)Zr_((1-x))Nb₂₄O_((52-62α)), M1_(x)Zn_((1-x))Nb₁₇O_((43.5-43.5α)) (i.e. Zn₂Nb₃₄O₈₇ base structure), M1_(x)Cu_((1-x))Nb₁₇O_((43.5-43.5α)) (i.e. Cu₂Nb₃₄O₈₇ base structure), M1_(x)W_((1-x))Nb_(0.571)O_((4.429-4.429α)) (i.e. W₇Nb₄O₃₁ base structure), M1_(x)W_((1-x))Nb_(0.889)O_((5.222-5.222α)) (i.e. W₉Nb₈O₄₇ base structure), M1_(x)W_((1-x))Nb_(3.2)O_((11-11α)) (i.e. WNb₁₆O₅₅ base structure), M1_(x)W_((1-x))Nb_(1.125)O_((5.813-5.813α)) (i.e. W₁₆Nb₁₈O₉₃ base structure), M1_(x)Al_((1-x))Nb₁₁O_((29-29α)), M1_(x)Ga_((1-x))Nb₁₁O_((29-29α)), M1_(x)Fe_((1-x))Nb₁₁O_((29-29α)), M1_(x)Al_((1-x))Nb₄₉O_((124-124α)), M1_(x)Ga_((1-x))Nb₄₉O_((124-124α)), M1_(x)Fe_((1-x))Nb₄₉O_((124-124α)), and M1_(x)Ge_((1-x))Nb₁₈O_((47-47α)) wherein α satisfies 0≤α≤0.05 and x and/or α is >0. x is as defined above. These represent modified versions of the ‘base’ mixed niobium oxide (i.e. when x=α=0). When x>0 the oxide is modified by cation substation of M1. When α>0 the oxide is modified by oxygen deficiency. Preferably the mixed niobium oxide is selected from M1_(x)Mo_((1-x))Nb₁₂O_((33-33α)), M1_(x)W_((1-x))Nb₁₂O_((33-33α)), M1_(x)V_((1-x))Nb₉O_((25-25α)), M1_(x)Zr_((1- x))Nb₂₄O_((62-62α)), M1_(x)Zn_((1-x))Nb₁₇O_((43.5-43.5α)) (i.e. Zn₂Nb₃₄O₈₇ base structure), M1_(x)Al_((1-x))Nb₁₁O_((29-29α)), M1_(x)W_((1- x))Nb_(0.571)O_((4.429-4.429α)) (i.e. W₇Nb₄O₃₁ base structure), and M1_(x)Ge_((1-x))Nb₁₈O_((47-47α)) wherein α satisfies 0≤α≤0.05 and x and/or α is >0. Most preferably, the mixed niobium oxide is selected from M1_(x)Mo_((1-x))Nb₁₂O_((33-33α)), M1_(x)W_((1-x))Nb₁₂O_((33-33α)), M1_(x)V_((1-x))Nb₉O_((25-25α)), M1_(x)Zr_((1-x))Nb₂₄O_((62-62α)), and M1_(x)W_((1-x))Nb_(0.571)O_((4.429-4.429α)) (i.e. W₇Nb₄O₃₁ base structure) wherein α satisfies 0≤α≤0.05 and x and/or α is >0.

It will be understood that the discussion of the variables of the active electrode material is intended to be read in combination. For example, M2 may be selected from one or more of Mo, W, V, Zr, P, Al, Zn, Ga, Ge, Ta, Cr, Cu, K, Mg, Ni, and Hf and M1 may be selected from P, B, Ti, Mg, V, Cr, W, Zr, Mo, Cu, Fe, Ga, Ge, K, Ni, Co, Al, Hf, Ta, and Zn. M2 may be selected from one or more of Mo, W, V, Zr, P, Al, Zn, Ga, and Ge and M1 may be selected from one or more of P, B, Ti, Mg, V, Cr, W, Zr, Mo, Ga, Ge, Al, and Zn. M2 may be selected from one or more of Mo, W, V, and Zr and M1 may be selected from one or more of Ti, Zr, V, W, and Mo. Optionally M1 and M2 are not Fe.

In one particular example M2 is selected from one or more of Mo, W, V, Zr, P, Al, Zn, Ga, Ge, Ta, Cr, Cu, K, Mg, Ni, and Hf; M1 is selected from P, B, Ti, Mg, V, Cr, W, Zr, Mo, Cu, Fe, Ga, Ge, K, Ni, Co, Al, Hf, Ta, and Zn; x satisfies 0.01<x<0.4. In a further example M2 is selected from one or more of Mo, W, V, Zr, P, Al, Zn, Ga, and Ge; M1 is selected from one or more of P, B, Ti, Mg, V, Cr, W, Zr, Mo, Ga, Ge, Al, and Zn; and 0.05≤x≤0.25. In a further example the mixed niobium oxide is selected from M1_(x)Mo_((1-x))Nb₁₂O_((33-33α)) and M1_(X)W_((1-x))Nb_(0.571)O_((4.429-4.429α)) (i.e. W₇Nb₄O₃₁ base structure) wherein α satisfies 0≤α≤0.05, x satisfies 0.01<x<0.4, wherein M1 is selected from Ti, Zr, V, W, and Mo.

The mixed niobium oxide may have a ReO₃-derived MO_(3-x) crystal structure. Preferably, the mixed niobium oxide has a Wadsley-Roth or Tetragonal Tungsten Bronze (“TTB” or “bronze”) crystal structure. Both Wadsley-Roth and bronze crystal structures are considered to be a crystallographic off-stoichiometry of the MO₃ (ReO₃) crystal structure, with simplified formula of MO_(3-x). As a result, these structures typically contain [MO_(B)] octahedral subunits in their crystal structure alongside others. Mixed niobium oxides with these structures are believed to have advantageous properties for use as active electrode materials, e.g. in lithium-ion batteries.

The open tunnel-like MO₃ crystal structure of MNOs also makes them ideal candidates for high capacity and high rate intercalation. The crystallographic off-stoichiometry that is introduced in MO_(3-x) structures causes crystallographic superstructures such as the Wadsley-Roth shear and the Bronze structures. These superstructures, compounded by other qualities such as the Jahn-Teller effect and crystallographic disorder by making use of multiple mixed cations, stabilise the crystal and keep the tunnels open and stable during intercalation, enabling extremely high rate performance.

The crystal formula of a charge balanced and thermodynamically stable Wadsley-Roth crystal structure obeys the following formula:

(M₁,M₂,M₃, . . . )_(mnp+1)O_(3mnp−(m+n)p+4)  (1)

In this formula, O is oxygen (the anion) and M (the cation) can be any alkali metal, alkali earth metal, transition element, semi-metal, or non-metal if the correct proportions are used to provide a stable structure. In the MNO, at least one of (M₁, M₂, M₃ . . . ) comprises Nb.

Formula (1) is based on crystal topography: m and n are the dimensions of the formed edge sharing superstructure blocks, ranging from 3-5 (integers). At the corner, blocks are connected into infinite ribbons (p=∞) only by edge-sharing, into pairs (p=2) by partly edge-sharing and partly tetrahedra or into isolated blocks only by tetrahedra (p=1). When p is infinity the formula becomes:

(M₁,M₂,M₃, . . . )_(mn)O_(3mn−(m+n))  (2)

More information can be found in work by Griffith et al.^([5])

Together, formula (1) and (2) define the full composition range for Wadsley-Roth crystal structures. The total crystal composition should also be charge neutral and thermodynamically favourable to follow the above description. Structures partially deficient in their oxygen content through introduction of oxygen vacancy defects are preferable when reducing the material's electrical resistance such that M_(x)O_(y) becomes M_(x)O_(γ-δ) where 0%<δ≤5%, i.e. the oxygen content is reduced by up to 5 atomic % relative to the amount of oxygen present.

Tetragonal tungsten bronze crystal structures are phases formed of a framework of [MO₆] octahedra sharing corners linked in such a way that three, four and five sided tunnels are formed (Montemayor et al.,^([6]) e.g. M₈W₉O₄₇). A bronze structure does not have to include tungsten^([7]). A number of 5-sided tunnels are filled with (M1, M2, M3 . . . ), 0, or a suitable cation to form the pentagonal columns. In the structure the pentagonal bipyramid MO₇ shares edge with five MO₆ octahedra. In the MNO, at least one of (M₁, M₂, M₃ . . . ) comprises Nb. Structures partially deficient in their oxygen content through introduction of oxygen vacancy defects are preferable when reducing the materials electrical resistance such that M_(x)O_(y) becomes M_(x)O_(y-δ) where 0%≤δ≤5%, i.e. the oxygen content is reduced by up to 5 atomic % relative to the amount of oxygen present.

The crystal structure of a material may be determined by analysis of X-ray diffraction (XRD) patterns, as is widely known. For instance, XRD patterns obtained from a given material can be compared to known XRD patterns to confirm the crystal structure, e.g. via public databases such as the ICDD (JCPDS) crystallography database. Rietveld analysis can also be used to determine the crystal structure. Therefore, the mixed niobium oxide may have a Wadsley-Roth or Tetragonal Tungsten Bronze crystal structure, as determined by X-ray diffraction.

Optionally, the crystal structure of the mixed niobium oxide, as determined by X-ray diffraction, corresponds to the crystal structure of the unmodified form of the mixed niobium oxide, wherein the unmodified form is expressed by the formula [M2][Nb]_(y)[O]z wherein M2 consists of a single element and wherein the unmodified form is not oxygen deficient, wherein the unmodified form is selected from one or more of: M2^(I)Nb₅O₁₃, M2^(I)BNb_(10.8)O₃₀, M2^(II)Nb₂O₆, M2^(II)2Nb₃₄O₃₇, M2^(III)Nb₁₁O₂₉, M2^(III)Nb₄₉O₁₂₄ (M2^(III) _(0.5)Nb_(24.5)O₆₂), M2^(IV)Nb₂₄O₆₂, M2^(IV)Nb₂O₇, M2^(IV) ₂Nb₁₀O₂₉, M2^(IV) ₂Nb₁₄O₃₉, M2^(IV)Nb₁₄O₃₇, M2^(IV)Nb₆O₁₇, M2^(IV)Nb₁₈O₄₇, M2^(V)Nb₉O₂₅, M2^(V) ₄Nb₁₈O₅₅, M2^(V) ₃Nb₁₇O₅₀, M2^(VI)Nb₁₂O₃₃, M2^(VI) ₄Nb₂₃O₇₇, M2^(VI) ₃Nb₁₄O₄₄, M2^(VI) ₅Nb₁₆O₅₅, M2^(VI) ₈Nb₁₈O₅₉, M2^(VI)Nb₂O₃, M2^(VI)Nb₁₈O₉₃, M2^(VI) ₂₀Nb₂₂O₁₁₅, M2^(VI) ₉Nb₈O₄₇, M2^(VI) ₈₂Nb₅₄O₃₈₁, M2^(VI) ₃₁Nb₂₀O₁₄₃, M2^(VI) ₇Nb₄O₃₁, M2^(VI) ₁₅Nb₂O₅₀, M2^(VI) ₃Nb₂O₁₄, and M2^(VI)Nb₁₂O₆₃, wherein the numerals I, II, III, IV, V, and VI represent the oxidation state of M2. In this way, it can be confirmed that the unmodified mixed niobium oxide has been modified without significantly affecting the crystal structure. Preferably the crystal structure of the mixed niobium oxide, as determined by X-ray diffraction, corresponds to the crystal structure of one or more of M2^(II) ₂Nb₃₄O₈₇, M2^(III)Nb₁₁O₂₉, M2^(III)Nb₄₉O₁₂₄, M2^(IV)Nb₂₄O₆₂, M2^(IV)Nb₁₈O₄₇, M2^(V)Nb₉O₂₅, M2^(VI)Nb₁₂O₃₃, M2^(VI) ₇Nb₄O₃₁, M2^(VI) ₉Nb₈O₄₇, M2^(VI) ₅Nb₁₆O₅₅, and M2^(VI) ₁₆Nb₁₈O₉₃.

The crystal structure of the mixed niobium oxide, as determined by X-ray diffraction, may correspond to the crystal structure of one or more of: MoNb₁₂O₃₃, WNb₁₂O₃₃, Mo₃Nb₁₄O₄₄, VNb₉O₂₅, ZrNb₂₄O₆₂, Zn₂Nb₃₄O₇, Cu₂Nb₃₄O₈₇, W₇Nb₄O₃₁, W₉Nb₈O₄₇, W₅Nb₁₆O₅₅, W₁₆Nb₁₈₀₉₃, AlNb₁₁O₂₉, GaNb₁₁O₂₉, FeNb₁₁O₂₉, AlNb₄O₁₂₄, GaNb₄₉O₁₂₄, FeNb₄₉O₁₂₄, and GeNb₁₃O₄₇. Preferably, the crystal structure of the mixed niobium oxide, as determined by X-ray diffraction, corresponds to the crystal structure of one or more of MoNb₁₂O₃₃, WNb₁₂O₃₃, ZrNb₂₄O₆₂, Zn₂Nb₃₄O₈₇, VNb₉O₂₅, W₅Nb₁₆O₅, AlNb₁₁O₂₉, GeNb₁₈O₄₇, W₇Nb₄O₃₁, and W₉Nb₈O₄₇. Most preferably the crystal structure of the mixed niobium oxide, as determined by X-ray diffraction, corresponds to the crystal structure of one or more of MoNb₁₂O₃₃, WNb₁₂O₃₃, ZrNb₂₄O₆₂, VNb₉O₂₅, and W₇Nb₄O₃₁.

Here the term ‘corresponds’ is intended to reflect that peaks in an X-ray diffraction pattern may be shifted by no more than 0.5 degrees (preferably shifted by no more than 0.25 degrees, more preferably shifted by no more than 0.1 degrees) from corresponding peaks in an X-ray diffraction pattern of the material listed above (e.g. M^(VI)Nb₁₂O₃₃ where M^(VI)=Mo etc.). This comparison may be performed with respect to the strongest peaks in the pattern, for example the three strongest peaks. Optionally, the crystal structure of the mixed niobium oxide does not correspond to the crystal structure of TiNb₂O₇, for example, optionally the measured XRD diffraction pattern of the mixed niobium oxide does not correspond to the JCPDS crystallography database entry database 00-039-1407, for TiNb₂O₇ Optionally, the crystal structure of the mixed niobium oxide does not correspond to the crystal structure of Ti₂Nb₁₀O₂₉. Optionally, the crystal structure of the mixed niobium oxides does not correspond to the crystal structure of M^(III)Nb₁₁O₂₉ for example FeNb₁₁O₂₉, GaNb₁₁O₂₉, CrNb₁₁O₂₉, and AlNb₁₁O₂₉.

The mixed niobium oxide and/or the lithium titanium oxide may further comprise Li and/or Na. For example, Li and/or Na may enter the crystal structures when the active electrode material is used in a metal-ion battery electrode.

The mixed niobium oxide may have a lithium diffusion rate of greater than 10⁻¹⁴ cm² s⁻¹. It may be advantageous to provide materials having a suitably high lithium diffusion rate, as this can provide improved performance in an electrochemical device comprising the active electrode material. For example, the lithium diffusion rate may be determined by cyclic voltammetry.

The specific capacity of the active electrode material may be 162 mAh/g or more. Here, specific capacity is defined as that measured in the 2nd cycle of a half cell galvanostatic cycling test at a rate of 0.1 C with a voltage window of 1.1-3.0V vs Li/Li+. It may be advantageous to provide materials having a high specific capacity, as this can provide improved performance in an electrochemical device comprising the active electrode material. The specific capacity may be targeted to a certain value by varying the proportion of the mixed niobium oxide and the lithium titanium oxide. Values of above 200 mAh/g can be achieved by using a high proportion of mixed niobium oxide, as shown by the present examples.

The mixed niobium oxide is preferably in particulate form. The mixed niobium oxide may have a Doo particle diameter in the range of 0.1-100 μm, or 0.5-50 μm, or 1-25 μm. These particle sizes are advantageous because they are easy to process and fabricate into electrodes. Moreover, these particle sizes avoid the need to use complex and/or expensive methods for providing nanosized particles. Nanosized particles (e.g. particles having a D₅₀ particle diameter of 100 nm or less) are typically more complex to synthesise and require additional safety considerations.

The mixed niobium oxide may have a D₁₀ particle diameter of at least 0.05 μm, or at least 0.1 μm, or at least 0.5 μm, or at least 1 μm. By maintaining a D10 particle diameter within these ranges, the potential for parasitic reactions in a Li ion cell is reduced from having reduced surface area, and it is easier to process with less binder in the electrode slurry.

The mixed niobium oxide may have a D₅₀ particle diameter of no more than 200 μm, no more than 100 μm, no more than 50 μm, or no more than 30 μm. By maintaining a D₅₀ particle diameter within these ranges, the proportion of the particle size distribution with large particle sizes is minimised, making the material easier to manufacture into a homogenous electrode.

The term “particle diameter” refers to the equivalent spherical diameter (esd), i.e. the diameter of a sphere having the same volume as a given particle, where the particle volume is understood to include the volume of any intra-particle pores. The terms “D_(n)” and “D_(n) particle diameter” refer to the diameter below which n % by volume of the particle population is found, i.e. the terms “D₅₀” and “D₅₀ particle diameter” refer to the volume-based median particle diameter below which 50% by volume of the particle population is found. Where a material comprises primary crystallites agglomerated into secondary particles, it will be understood that the particle diameter refers to the diameter of the secondary particles. Particle diameters can be determined by laser diffraction. For example, particle diameters can be determined in accordance with ISO 13320:2009.

The lithium titanium oxide is in preferably in particulate form. The lithium titanium oxide may have a D₅₀ particle diameter in the range of 0.1-50 μm, or 0.25-20 μm, or 0.5-15 μm. The lithium titanium oxide may have a D₁₀ particle diameter of at least 0.01 μm, or at least 0.1 μm, or at least 0.5 μm. The lithium titanium oxide may have a D₉₀ particle diameter of no more than 100 μm, no more than 50 μm, or no more than 25 μm. By maintaining a D₉₀ particle diameter in this range the packing of lithium titanium oxide particles in the mixture with mixed niobium oxide particles is improved.

Lithium titanium oxides are typically used in battery anodes at small particle sizes due to the low electronic conductivity of the material. In contrast, the mixed niobium oxide as defined herein may be used at larger particle sizes since it typically has a higher lithium ion diffusion coefficient than lithium titanium oxide. Advantageously, in the active electrode material the lithium titanium oxide may have a smaller particle size than the mixed niobium oxide, for example such that the ratio of the D₅₀ particle diameter of the lithium titanium oxide to the D₅₀ particle diameter of the mixed niobium oxide is in the range of 0.01:1 to 0.9:1, or 0.1:1 to 0.7:1. In this way, the smaller lithium titanium oxide particles may be accommodated in the voids between the larger mixed niobium oxide particles, increasing the packing efficiency of the active electrode material.

The mixed niobium oxide may have a BET surface area in the range of 0.1-100 m²/g, or 0.5-50 m²/g, or 1-20 m²/g. The lithium titanium oxide may have a BET surface area in the range of 0.1-100 m²/g, or 1-50 m²/g, or 3-30 m²/g. In general, a low BET surface area is preferred in order to minimise the reaction of the active electrode material with the electrolyte, e.g. minimising the formation of solid electrolyte interphase (SEI) layers during the first charge-discharge cycle of an electrode comprising the material. However, a BET surface area which is too low results in unacceptably low charging rate and capacity due to the inaccessibility of the bulk of the active electrode material to metal ions in the surrounding electrolyte. The the ratio of the BET surface area of the lithium titanium oxide to the BET surface area of the mixed niobium oxide is in the range of 1.1:1 to 20:1, or 1.5:1 to 10:1.

The term “BET surface area” refers to the surface area per unit mass calculated from a measurement of the physical adsorption of gas molecules on a solid surface, using the Brunauer-Emmett-Teller theory. For example, BET surface areas can be determined in accordance with ISO 9277:2010.

The mixed niobium oxide may comprise a carbon coating. The coating may be present in an amount of up to 10 wt %, or 0.05-5 wt %, or 0.1-3 wt %, based on the total weight of the mixed niobium oxide and the coating. It has been found that a carbon precursor comprising polyaromatic sp² carbon provides a particularly beneficial carbon coating on mixed niobium oxides. Preferably the carbon coating comprises polyaromatic sp² carbon. Such a coating is formed by pyrolysing a carbon precursor comprising polyaromatic sp² carbon since the sp² hybridisation is largely retained during pyrolysis. Typically, pyrolysis of a polyaromatic sp² carbon precursor under reducing conditions results in the domains of sp² aromatic carbon increasing in size. Accordingly, the presence of a carbon coating comprising polyaromatic sp² may be established via knowledge of the precursor used to make the coating. The carbon coating may be defined as a carbon coating formed from pyrolysis of a carbon precursor comprising polyaromatic sp² carbon. Preferably, the carbon coating is derived from pitch carbons.

The presence of a carbon coating comprising polyaromatic sp² carbon may also be established by routine spectroscopic techniques. For instance, Raman spectroscopy provides characteristic peaks (most observed in the region 1,000-3,500 cm¹) which can be used to identify the presence of different forms of carbon. A highly crystalline sample of sp³ carbon (e.g. diamond) provides a narrow characteristic peak at ˜1332 cm⁻¹. Polyaromatic sp² carbon typically provides characteristic D, G, and 2D peaks. The relative intensity of D and G peaks (I_(D)/I_(G)) can provide information on the relative proportion of sp² to sp³ carbon. The mixed niobium oxide may have an I_(D)/I_(G) ratio as observed by Raman spectroscopy within the range of 0.85-1.15, or 0.90-1.10, or 0.95-1.05.

X-ray diffraction may also be used to provide information on the type of carbon coating. For example, an XRD pattern of a mixed niobium oxide with a carbon coating may be compared to an XRD pattern of the uncoated mixed niobium oxide and/or to an XRD pattern of a pyrolysed sample of the carbon precursor used to make the carbon coating.

The carbon coating may be semi-crystalline. For example, the carbon coating may provide a peak in an XRD pattern of the mixed niobium oxide centred at 2θ of about 26° with a width (full width at half maximum) of at least 0.20°, or at least 0.25°, or at least 0.30°.

The lithium titanium oxide preferably has a spinel or ramsdellite crystal structure, e.g. as determined by X-ray diffraction. An example of a lithium titanium oxide having a spinel crystal structure is Li₄Ti₅O₁₂. An example of a lithium titanium oxide having a ramsdellite crystal structure is Li₂Ti₃O₇. These materials have been shown to have good properties for use as active electrode materials. Therefore, the lithium titanium oxide may have a crystal structure as determined by X-ray diffraction corresponding to Li₄Ti₅O₁₂ and/or Li₂Ti₃O₇. The lithium titanium oxide may be selected from Li₄Ti₅O₁₂, Li₂Ti₃O₇, and mixtures thereof.

The lithium titanium oxide may be doped with additional cations or anions. The lithium titanium oxide may be oxygen deficient. The lithium titanium oxide may comprise a coating, optionally wherein the coating is selected from carbon, polymers, metals, metal oxides, metalloids, phosphates, and fluorides.

The lithium titanium oxide may be synthesises by conventional ceramic techniques, for example solid-state synthesis or sol-gel synthesis. Alternatively, the lithium titanium oxide may be obtained from a commercial supplier.

A method of making a mixed niobium oxide for use in the invention comprises the steps of: providing one or more precursor materials; mixing said precursor materials to form a precursor material mixture; and heat treating the precursor material mixture in a temperature range from 400° C.-1350° C. to form the mixed niobium oxide.

The one or more precursor materials may include an M1 source, an M2 source, and a source of Nb. It will be understood that the sources may be contaminated by impurities. For example, Ta is a typical impurity present in sources of Nb which may thus be present in a mixed niobium oxide.

The phrase ‘M1 source’ is used herein to describe a material comprising M1 ions/atoms. The phrase ‘M2 source’ is used herein to describe a material comprising M2 ions/atoms. The phrase ‘a source of Nb’ is used herein to describe a material comprising Nb ions/atoms, as appropriate.

The precursor materials may include one or more metal oxides, metal hydroxides, metal salts or oxalates. For example, the precursor materials may include one or more metal oxides of different oxidation states and/or of different crystal structure. Examples of suitable metal oxide precursor materials include but are not limited to: Nb₂O₅, NbO₂, WO₃, TiO₂, MoO₃, V₂O₅, ZrO₂, and MgO. However, the precursor materials may not comprise a metal oxide, or may comprise ion sources other than oxides. For example, the precursor materials may comprise metal salts (e.g. NO₃ ⁻, SO₃ ⁻) or other compounds (e.g. oxalates).

Some or all of the precursor materials may be particulate materials. Where they are particulate materials, preferably they have D₅₀ particle diameter of <20 μm in diameter. The D₅₀ particle diameter may be in a range from e.g. 10 nm to 20 μm. Providing particulate materials with such a particle size can help to promote more intimate mixing of precursor materials, thereby resulting in more efficient solid-state reaction during the heat treatment step. However, it is not essential that the precursor materials have an initial D₅₀ particle diameter of <20 μm, as the particle size of the one or more precursor materials may be mechanically reduced during the step of mixing said precursor materials to form a precursor material mixture.

The step of mixing/milling the precursor materials to form a precursor material mixture may be performed by a process selected from (but not limited to): dry or wet planetary ball milling, rolling ball milling, high shear milling, airjet milling, and/or impact milling. The force used for mixing/milling may depend on the morphology of the precursor materials. For example, where some or all of the precursor materials have larger particle sizes (e.g. a D₅₀ particle diameter of greater than 20 μm), the milling force may be selected to reduce the particle size of the precursor materials such that the such that the D₅₀ particle diameter of the precursor material mixture is reduced to 20 μm or lower. When the D₅₀ particle diameter of particles in the precursor material mixture is 20 μm or less, this can promote a more efficient solid-state reaction of the precursor materials in the precursor material mixture during the heat treatment step.

The step of heat treating the precursor material mixture may be performed for a time of from 1 hour to 24 hours, more preferably from 3 hours to 14 hours. For example, the heat treatment step may be performed for 1 hour or more, 2 hours or more, 3 hours or more, 6 hours or more, or 12 hours or more. The heat treatment step may be performed for 24 hours or less, 18 hours or less, 14 hours or less, or 12 hours or less.

In some methods it may be beneficial to perform a two-step heat treatment. For example, the precursor material mixture may be heated at a first temperature for a first length of time, follow by heating at a second temperature for a second length of time. The second temperature may be higher than the first temperature. Performing such a two-step heat treatment may assist the solid state reaction to form the desired crystal structure.

The step of heat treating the precursor material mixture may be performed in a gaseous atmosphere. The gaseous atmosphere may be an inert atmosphere, or may be a reducing atmosphere. Where it is desired to make an oxygen-deficient material, preferably the step of heat treating the precursor material mixture is performed in an inert or reducing atmosphere. Suitable gaseous atmospheres comprise: air, N₂, Ar, He, CO₂, CO, O₂, H₂, and mixtures thereof.

The method may include one or more post-processing steps after formation of the mixed niobium oxide.

In some cases, the method may include a post-processing step of heat treating the mixed niobium oxide, sometimes referred to as ‘annealing’. This post-processing heat treatment step may be performed in a different gaseous atmosphere to the step of heat treating the precursor material mixture to form the mixed niobium oxide. The post-processing heat treatment step may be performed in an inert or reducing gaseous atmosphere. Such a post-processing heat treatment step may be performed at temperatures of above 500° C., for example at about 900° C. Inclusion of a post-processing heat treatment step may be beneficial to e.g. form deficiencies or defects in the mixed niobium oxide, for example to form oxygen deficiencies.

In some cases, the method may include a post-processing step of mixing the mixed niobium oxide with a carbon source, and thereby forming a carbon coating on the mixed niobium oxide. Optionally, the mixture of the mixed niobium oxide and the carbon source may be heated to thereby form the carbon coating on the mixed niobium oxide. Suitable carbon sources include but are not limited to: carbohydrate materials (e.g. sugars, polymers); conductive carbons (e.g. carbon black); and/or aromatic carbon materials (e.g. pitch carbon).

One method of forming a carbon coating includes a step of milling the mixed niobium oxide with a carbon source, followed by pyrolysis of the mixed niobium oxide and carbon source (e.g. in a furnace) under an inert or reducing atmosphere.

Another preferred method of forming a carbon coating includes mixing of the mixed niobium oxide with a carbon source, dispersion of the mixed niobium oxide and carbon source in an aqueous slurry, followed by spray drying. The resulting powder may optionally be pyrolysed. Where the carbon source is e.g. conductive carbon black, it is not necessary to pyrolyse the material post spray-drying.

In some cases, the method may include a post-processing step of milling the mixed niobium oxide to modify the mixed niobium oxide particle size. For example, the mixed niobium oxide may be treated by one or more processes including air jet milling, impact milling, high shear milling, sieving, or ball milling. This may provide a more suitable particle size for use in desired applications of the mixed niobium oxide.

It has been found that a carbon precursor comprising polyaromatic sp² carbon provides a particularly beneficial carbon coating on mixed niobium oxides for use in the invention. Therefore, a method of making a coated mixed niobium oxide may comprise the steps of: combining a mixed niobium oxide with a carbon precursor comprising polyaromatic sp² carbon to form an intermediate material; and heating the intermediate material under reducing conditions to pyrolyse the carbon precursor forming a carbon coating on the mixed niobium oxide and introducing oxygen vacancies into the mixed niobium oxide.

The intermediate material may comprise the carbon precursor in an amount of up to 25 wt %, or 0.1-15 wt %, or 0.2-8 wt %, based on the total weight of the mixed niobium oxide and the carbon precursor. The carbon coating on the mixed niobium oxide may be present in an amount of up to 10 wt %, or 0.05-5 wt %, or 0.1-3 wt %, based on the total weight of the mixed niobium oxide and coating. These amounts of the carbon precursor and/or carbon coating provide a good balance between improving the electronic conductivity by the carbon coating without overly reducing the capacity of the mixed niobium oxide by overly reducing the proportion of the mixed niobium oxide. The mass of carbon precursor lost during pyrolysis may be in the range of 30-70 wt %.

The step of heating the intermediate material under reducing conditions may be performed at a temperature in the range of 400-1,200° C., or 500-1,100° C., or 600-900° C. The step of heating the intermediate material under reducing conditions may be performed for a duration within the range of 30 minutes to 12 hours, 1-9 hours, or 2-6 hours.

The step of heating the intermediate material under reducing conditions may be performed under an inert gas such as nitrogen, helium, argon; or may be performed under a mixture of an inert gas and hydrogen; or may be performed under vacuum.

The carbon precursor comprising polyaromatic sp² carbon may be selected from pitch carbons, graphene oxide, graphene, and mixtures thereof. Preferably, the carbon precursor comprising polyaromatic sp² carbon is selected from pitch carbons, graphene oxide, and mixtures thereof. Most preferably, the carbon precursor comprising polyaromatic sp² carbon is selected from pitch carbons. The pitch carbons may be selected from coal tar pitch, petroleum pitch, mesophase pitch, wood tar pitch, isotropic pitch, bitumen, and mixtures thereof.

Pitch carbon is a mixture of aromatic hydrocarbons of different molecular weights. Pitch carbon is a low cost by-product from petroleum refineries and is widely available. The use of pitch carbon is advantageous because pitch has a low content of oxygen. Therefore, in combination with heating the intermediate material under reducing conditions, the use of pitch favours the formation of oxygen vacancies in the mixed niobium oxide.

Other carbon precursors typically contain substantial amounts of oxygen. For example, carbohydrates such as glucose and sucrose are often used as carbon precursors. These have the empirical formula C_(m)(H₂O)_(n) and thus contain a significant amount of covalently-bonded oxygen (e.g. sucrose has the formula C₁₂H₂₂O₁₁ and is about 42 wt % oxygen). In some instances the pyrolysis of carbon precursors which contain substantial amounts of oxygen may prevent or inhibit reduction of a mixed niobium oxide, or even lead to oxidation, meaning that oxygen vacancies may not be introduced into the mixed niobium oxide. Accordingly, the carbon precursor may have an oxygen content of less than 10 wt %, preferably less than 5 wt %.

The carbon precursor may be substantially free of sp³ carbon. For example, the carbon precursor may comprise less than 10 wt % sources of sp³ carbon, preferably less than 5 wt % sources of sp³ carbon. Carbohydrates are sources of sp³ carbon. The carbon precursor may be free of carbohydrates. It will be understood that some carbon precursors used may contain impurities of sp³ carbon, for example up to 3 wt %.

The invention also provides a composition comprising the active electrode material of the first aspect of the invention and at least one other component, optionally wherein the at least one other component is selected from a binder, a solvent, a conductive additive, an additional active electrode material, and mixtures thereof. Such a composition is useful for preparing an electrode, e.g. an anode for a lithium-ion battery.

The invention also provides an electrode comprising the active electrode material of the first aspect of the invention in electrical contact with a current collector. The electrode may form part of a cell. The electrode may form an anode as part of a lithium-ion battery. Preferably, the active electrode material is in the form of an active layer on the current collector, wherein the active layer has a density of 2.00-3.75 g cm⁻³.. It may be advantageous to provide materials having such an electrode density, as this can provide improved performance in an electrochemical device comprising the active electrode material. Specifically, when the electrode density is high, high volumetric capacities can be achieved, as gravimetric capacity x electrode density x active electrode material fraction=volumetric capacity.

The invention also provides the use of the active electrode material of the first aspect of the invention in an anode for a metal-ion battery, optionally wherein the metal-ion battery is a lithium-ion battery.

A further implementation of the invention is an electrochemical device comprising an anode, a cathode, and an electrolyte disposed between the anode and the cathode, wherein the anode comprises an active electrode material according to the first aspect of the invention; optionally wherein the electrochemical device is a lithium-ion battery or a sodium-ion battery. Preferably, the electrochemical device is a lithium-ion battery having a reversible anode active material specific capacity of greater than 165 mAh/g at 20 mA/g, wherein the battery can be charged and discharged at current densities relative to the anode active material of 200 mA/g or more, or 1000 mA/g or more, or 2000 mA/g or more, or 4000 mA/g or more whilst retaining greater than 70% of the initial cell capacity at 20 mA/g. It has been found that use of the active electrode materials of the first aspect of the invention can enable the production of a lithium-ion battery with this combination of properties, representing a lithium-ion battery that is particularly suitable for use in applications where high charge and discharge current densities are desired. Notably, the examples have shown that active electrode materials according to the first aspect of the invention have excellent capacity retention at high C-rates.

Preferably, the electrochemical device is a lithium-ion battery cell. The anode active material mixture in the cell preferably having an initial coulombic efficiency greater than 88% or greater than 90%. Initial coulombic efficiency has been measured as the difference in the lithiation and de-lithiation capacity on the 1^(st) charge/discharge cycle at C/10 in a half-cell. It may be advantageous to provide materials having a suitably high initial coulombic efficiency, as this can provide improved performance in an electrochemical device comprising the active electrode material.

In the second aspect, the invention provides a method for making an active electrode material, wherein the active electrode material is as defined in the first aspect, the method comprising mixing at least one lithium titanium oxide with at least one mixed niobium oxide. The lithium titanium oxide and the mixed niobium oxide are as defined above.

The step of mixing at least one lithium titanium oxide with at least one mixed niobium oxide may comprise low to high energy powder mixing/blending techniques, such as rotational mixing in multiple directions, rotational V-type blending over a single axis, planetary mixing, centrifugal planetary mixing, and high shear mixing.

The step of mixing at least one lithium titanium oxide with at least one mixed niobium oxide may comprise mixing in a carrier solvent.

Prior to mixing, the method may include the step of milling and/or classifying the lithium titanium oxide, e.g. to provide any of the particle size parameters given above. Prior to mixing, the method may include the step of milling and/or classifying the mixed niobium oxide, e.g. to provide any of the particle size parameters given above. The method may include a step of milling and/or classifying the mixture of the lithium titanium oxide and the mixed niobium oxide. The milling and/or classifying may be performed by impact milling or jet milling.

Optionally, the method for making an active electrode material, wherein the active electrode material is as defined in the first aspect, comprises the steps of: combining a mixed niobium oxide with a carbon precursor comprising polyaromatic sp² carbon to form an intermediate material; heating the intermediate material under reducing conditions to pyrolyse the carbon precursor forming a carbon coating on the mixed niobium oxide and introducing oxygen vacancies into the mixed niobium oxide; and mixing at least one lithium titanium oxide with the coated mixed niobium oxide.

REFERENCE EXAMPLES

The following reference examples demonstrate the improvement in properties of a modified mixed niobium oxide (i.e. a cation substituted and/or oxygen deficient oxide) compared to the unmodified ‘base’ mixed niobium oxide. The reference examples test the oxides as the sole active electrode material. It would be expected that the same improvements would be seen when the oxides are tested in combination with a lithium titanium oxide in accordance with the invention, i.e. that a mixture of a modified mixed niobium oxide and a lithium titanium oxide will have improved properties for use as an active electrode material compared to a mixture of the unmodified ‘base’ mixed niobium oxide and the lithium titanium oxide.

A number of different materials were prepared and characterised, as summarised in Table 1, below. Broadly, these samples can be split into a number of groups:

Samples 1, 2, 3, 4, 5, 14, 15, 16, 18, and 22 belong to the same family of Wadsley-Roth phases based on MoNb₁₂O₃₃. Sample 1 is the base crystal structure, which is modified to a mixed metal cation structure by exchanging one or multiple cations in samples 2 to 4, and/or in a mixed crystal configuration (blending with isostructural WNb₁₂O₃₃) in samples 14, 15, 16, 18, and 22. Oxygen deficiencies are created in the base crystal in sample 5 and in the mixed metal cation structure 18. Sample 3 is a spray-dried and carbon-coated version of the crystal made in sample 2, and sample 22 is a spray-dried and carbon-coated version of the crystal made in sample 16.

Samples 6, 7, 17, 19, 20 belong to the same family of Wadsley-Roth phases based on ZrNb₂₄O₆₂ (M⁴⁺Nb₂₄O₆₂, 3×4 block of octahedra with half a tetrahedron at each block corner).

Samples 8, 9 and R11 belong to the same family of Wadsley-Roth phases based on WNb₁₂O₃₃ (M⁶⁺Nb₁₂O₃₃, a 3×4 NbO₆ octahedra block with a tetrahedron at each block corner).

Samples 10, 11 and 21 belong to the same family of Wadsley-Roth phases based on VNb₉₀₂₅ (M⁵⁺Nb₉O₂₅, a 3×3 NbO_(B) octahedra block with a tetrahedron at each block corner).

Samples 12, 13 and R14 belong to the same family of tungsten tetragonal bronzes (TTB) based on W₇Nb₄O₃₁ (M⁶*7Nb₄O₃₁). This is a tetragonal tungsten bronze structure, where MO₆ (M=0.4 Nb+0.6 W) octahedra are exclusively corner-sharing, with 3, 4, and 5-sided tunnels. Some of these tunnels are filled with —O-M-O— chains whereas others are open for lithium ion transport and storage.

Samples R1, R2, R13 belong to the same family of Wadsley-Roth phases based on Zn₂Nb₃₄O₈₇ (M²⁺2Nb₃₄O₈₇). This orthorhombic phase consists out of 3×4 blocks of MO_(B) octahedra (M=Zn⁺²/Nb⁺⁵), where the blocks are connected exclusively by edge-sharing and have no tetrahedra.

Samples R3, R4, R5, R12 belong to the same family of Wadsley-Roth phases based on AlNb₁₁O₂₉ (M³⁺Nb₁₁O₂₉). The structure belongs to monoclinic shear structure with 3×4 octahedra blocks connected through exclusively edge-sharing and have no tetrahedra.

Samples R6, R7, R8 belong to the same family of Wadsley-Roth phases based on GeNb₁₈O₄₇ (M⁴⁺Nb₁₃O₄₇). The structure is similar to sample 10 with 3×3 NbO₆ octahedra blocks and one tetrahedron connecting blocks at corners. However, the structure contains intrinsic defects due to Ge⁺⁴ instead of V⁵⁺.

Samples R9, R10 belong to the same family of Wadsley-Roth phases based on W₅Nb₁₆O₅₅ (M⁶⁺ ₅Nb₁₆O₅₅). The structure is made of 4×5 blocks connected at the sides by edge-sharing (W,Nb)O₆ and connected at the corners by WO₄ tetrahedra. This structure is similar to Sample 8 and 9 but with a larger block size.

TABLE 1 A summary of different compositions synthesised. Sample No. Composition Material Synthesis  1 * MoNb₁₂O₃₃ Solid state  2 Ti_(0.05)Mo_(0.95)Nb₁₂O₃₃ Solid state  3 Ti_(0.05)Mo_(0.95)Nb₁₂O₃₃ + C Solid state, spray dry, carbon pyrolysis  4 Zr_(0.05)Mo_(0.95)Nb₁₂O₃₃ Solid state  5 MoNb₁₂O_(33−δ) Solid state  6 * ZrNb₂₄O₆₂ Solid state  7 V_(0.05)Zr_(0.95)Nb₂₄O₆₂ Solid state  8 * WNb₁₂O₃₃ Solid state  9 Ti_(0.05)W_(0.95)Nb₁₂O₃₃ Solid state 10 * VNb₉O₂₅ Solid state 11 Ti_(0.05)V_(0.95)Nb₉O₂₅ Solid state 12 * W₇Nb₄O₃₁ (WNb_(0.57)O_(4.43)) Solid state 13 Ti_(0.05)W_(0.95)Nb_(0.57)O_(4.43) Solid state (Ti_(0.35)W_(6.65)Nb₄O₃₁) 14 W_(0.25)Mo_(0.75)Nb₁₂O₃₃ Solid state 15 Ti_(0.05)W_(0.25)Mo_(0.70)Nb₁₂O₃₃ Solid state 16 T_(i0.05)Zr_(0.05)W_(0.25)Mo_(0.65)Nb₁₂O₃₃ Solid state 17 Ti_(0.05)Zr_(0.95)Nb₂₄O₆₂ Solid state 18 Ti_(0.05)Zr_(0.05)W_(0.25)Mo_(0.65)Nb₁₂O_(33−δ) Solid state 19 Mo_(0.05)Zr_(0.95)Nb₂₄O₆₂ Solid state 20 Mo_(0.05)V_(0.05)Zr_(0.95)Nb₂₄O₆₂ Solid state 21 Mo_(0.05)V_(0.95)Nb₉O₂₅ Solid state 22 Ti_(0.05)Zr_(0.05)W_(0.25)Mo_(0.65)Nb₁₂O₃₃ + C Solid state, spray dry, carbon pyrolysis R1 * Zn₂Nb₃₄O₈₇ Solid state R2 Ge_(0.1)Zn_(1.9)Nb₃₄O₈₇ Solid state R3 * AlNb₁₁O₂₉ Solid state R4 Fe_(0.05)Al_(0.95)Nb₁₁O₂₉ Solid state R5 Ga_(0.05)Al_(0.95)Nb₁₁O₂₉ Solid state R6 * GeNb₁₈O₄₇ Solid state R7 K_(0.02)Co_(0.02)Ge_(0.96)Nb₁₈O₄₇ Solid state R8 K_(0.02)Co_(0.02)Ge_(0.96)Nb₁₈O_(47−α) Solid state R9 * W₅Nb₁₆O₅₅ Solid state R10 W₅Nb₁₆O_(55−α) Solid state R11 WNb₁₂O_(33−α) Solid state R12 AlNb₁₁O_(29−α) Solid state R13 Zn₂Nb₃₄O_(87−α) Solid state R14 W₇Nb₄O_(31−α) Solid state Samples indicated with * are comparative samples.

Material Synthesis

Samples listed in Table 1 were synthesised using a solid-state route. In a first step, metal oxide precursor commercial powders (Nb₂O₅, NbO₂, MoO₃, ZrO₂, TiO₂, WO₃, V₂O₅, ZrO₂, K₂O, CoO, Fe₂O₃, GeO₂, Ga₂O₃, Al₂O₃, ZnO and/or MgO) were mixed in stochiometric proportions and planetary ball-milled at 550 rpm for 3h in a zirconia jar and milling media with a ball to powder ratio of 10:1. The resulting powders were then heated in a static muffle furnace in air in order to form the desired crystal phase. Samples 1 to 5 and 12 to 16, 18 and 22 were heat-treated at 90000 for 12h; samples 6 to 9, 17, 19, and 20 were heat-treated at 1200° C. for 12h, with samples 6, 7, 17, 19 and 20 undergoing a further heat treatment step at 1350° C. for an additional 4h; samples 10, 11 and 21 were heat-treated at 1000° C. for 12h. Sample 3 and 22 were further mixed with a carbohydrate precursor (such as sucrose, maltodextrin or other water-soluble carbohydrates), dispersed in an aqueous slurry at concentrations of 5, 10, 15, or 20 w/w % with ionic surfactant, and spray-dried in a lab-scale spray-drier (inlet temperature 220° C., outlet temperature 95° C., 500 mL/h sample introduction rate). The resulting powder was pyrolyzed at 600° C. for 5h in nitrogen. Sample 5 and 18 were further annealed in nitrogen at 900° C. for 4 hours.

Samples R1, R2, R6, R7, R8, R9, R10 were prepared by ball milling as above, and impact milling at 20,000 rpm as needed to a particle size distribution with D90<20 μm, then heat-treated as in a muffle furnace in air at 1200° C. for 12 h; samples R8, R10, R11, R12, R13 were further annealed in nitrogen at 1000° C. for 4 h; R14 was annealed in nitrogen at 900° C. for 5 h. Samples R3, R4, R5 were heat-treated at 1300° C. for 12 h. Samples R1-R10 were de-agglomerated after synthesis by impact milling or jet milling to the desired particle size ranges.

XRD Characterisation of Samples

The phase purity of some samples was analysed using Rigaku Miniflex powder X-ray diffractometer in 2θ range (10-70°) at 1°/min scan rate.

FIG. 1 shows the measured XRD diffraction patterns for samples 1, 4, 14, 2, 5, 15, 16, 18, 22 which are relevant to Comparative Study A. All diffraction patterns have peaks at the same locations (within instrument error, that is 0.1°), and match JCPDS crystallography database entry database JCPDS 73-1322, which corresponds to MoNb₁₂O₃₃. There is no amorphous background noise and the peaks are sharp and intense. This means that all samples are phase-pure and crystalline, with crystal structure matching MoNb₁₂O₃₃.

FIG. 2 shows the measured XRD diffraction patterns for samples 8 and 9. All diffraction patterns have peaks at the same locations (within instrument error, that is 0.1°), and match JCPDS crystallography database entry database JCPDS 73-1322, which corresponds to WNb₁₂O₃₃. There is no amorphous background noise and the peaks are sharp and intense. This means that all samples are phase-pure and crystalline, with crystal structure matching WNb₁₂O₃₃.

FIG. 3 shows the measured XRD diffraction patterns for samples 6, 7, 17, 19, 20 which are relevant to Comparative Study B. All diffraction patterns have peaks at the same locations (within instrument error, that is 0.1°), and match JCPDS crystallography database entry database JCPDS 01-072-1655, which corresponds to ZrNb₂₄O₆₂. There is no amorphous background noise and the peaks are sharp and intense. This means that all samples are phase-pure and crystalline, with crystal structure matching ZrNb₂₄O₆₂.

FIG. 4 shows the measured XRD diffraction patterns for samples 10, 11, 21. All diffraction patterns have peaks at the same locations (within instrument error, that is 0.1°), and match JCPDS crystallography database entry database JCPDS 00-049-0289, which corresponds to VNb₉O₂₅. There is no amorphous background noise and the peaks are sharp and intense. This means that all samples are phase-pure and crystalline, with crystal structure matching VNb₉O₂₅.

FIG. 5 shows the measured XRD diffraction patterns for samples 12 and 13. All diffraction patterns have peaks at the same locations (within instrument error, that is 0.1°), and match JCPDS crystallography database entry database JCPDS 00-020-1320, which corresponds to W₇Nb₄O₃₁. There is no amorphous background noise and the peaks are sharp and intense. This means that all samples are phase-pure and crystalline, with crystal structure matching W₇Nb₄O₃₁.

FIG. R1 shows the measured XRD diffraction patterns for samples R1, R2. FIG. R10 shows the XRD pattern for sample R13. All diffraction patterns have peaks at the same locations (within 0.1-0.2°), and match JCPDS crystallography database entry JCPDS 22-353. There is no amorphous background noise and the peaks are sharp and intense. This means that all samples are phase-pure and crystalline, with crystallite size 52±12 nm according to the Scherrer equation and crystal structure matching Zn₂Nb₃₄O₈₇.

FIG. R2 shows the measured XRD diffraction patterns for samples R3, R4, R5. FIG. R10 shows the XRD pattern for sample R12. All diffraction patterns have peaks at the same locations (within 0.1-0.2°), and match JCPDS crystallography database entry JCPDS 72-159 (isostructural Ti₂Nb₁₀O₂₉). There is no amorphous background noise and the peaks are sharp and intense. This means that all samples are phase-pure and crystalline, with crystallite size 53±16 nm according to the Scherrer equation and crystal structure matching AlNb₁₁O₂₉.

FIG. R3 shows the measured XRD diffraction patterns for samples R6, R7, R8. All diffraction patterns have peaks at the same locations (within 0.1-0.2°), and match ICSD crystallography database entry 72683 (isostructural PNb₉O₂₅). There is no amorphous background noise and the peaks are sharp and intense. This means that all samples are phase-pure and crystalline, with crystallite size 53±3 nm according to the Scherrer equation and crystal structure matching GeNb₁₈O₄₇.

FIG. R4 shows the measured XRD diffraction patterns for samples R9, R10. All diffraction patterns have peaks at the same locations (within 0.1-0.2°), and match JCPDS crystallography database entry JCPDS 44-0467. There is no amorphous background noise and the peaks are sharp and intense. This means that all samples are phase-pure and crystalline, with crystallite size 37±11 nm according to the Scherrer equation and crystal structure matching WNb₁₅O₅₅.

TGA Characterisation of Samples

Thermogravimetric Analysis (TGA) was performed on some samples using a Perkin Elmer Pyris 1 system in a synthetic air atmosphere. Samples were first held for 15 min at 30° C., then heated from 30° C. to 950° C. at 5° C./min, and finally held for 30 min at 950° C. TGA was performed on sample 3 to quantify carbon content.

FIG. 6 shows TGA characterisation in air of sample 3. The sharp drop in mass between ˜400° C. and 500° C. is attributed to the decomposition of the carbon coating. The decomposition temperature corresponds to a mixture of amorphous and graphitic carbon. The amount of mass loss indicates that sample 3 includes 1.1 w. % of carbon coating, which is in line with the amount expected from the stoichiometry of the precursors.

Qualitative Assessment of Oxygen Deficiency

As discussed above, sample 5 and 18 were heat-treated at 900° C. for 12h to form the mixed niobium oxide, and was then further annealed in nitrogen (a reducing atmosphere) at 900° C., in a post-processing heat treatment step. A colour change from white to dark purple was observed after the post-processing heat treatment in nitrogen, indicating change in oxidation states and band structure of the material, as a result of oxygen deficiency of the sample.

Samples R8, R10, R11, R12, R13 were further annealed in nitrogen at 1000° C. for 4 h, sample R14 was annealed in nitrogen at 900° C. for 5 h. Sample R7 transitions from a white colour to a deep yellow colour upon introduction of induced oxygen deficiencies in sample R8; sample R9 transitions from an off-white colour to a blue-grey colour upon introduction of induced oxygen deficiencies in sample R10; sample 8 transitions from off-white to light blue in R11; sample R3 transitions from white to grey/black in R12; sample R1 transitions from white to grey/black in R3; sample 12 transitions from light yellow to dark blue in R14.

Particle Size Distribution Analysis of Samples

Particle Size Distributions were obtained with a Horiba laser diffraction particle analyser for dry powder. Air pressure was kept at 0.3 MPa. The results are set out in Table 2, below.

TABLE 2 Summary of particle size distribution statistics for samples 1, 2, 15, 16, 18, 3 before pyrolysis, 3 after pyrolysis, 16 and 18 after post-processing, and samples R1-R14. Sample D₁₀ [μm] D₅₀ [μm] D₉₀ [μm]  1* 3.8 11.2 50.0  2 2.6 10.9 87.4 15 3.6 21.2 55.3 16 4.7 31.2 82.9 18 5.1 57.7 176 3 before pyrolysis 4.2 8.2 16.3 3 after pyrolysis 6.7 12.7 51.1 16 after impaction 1.0 2.6 4.8 milling 18 after impaction 1.4 4.4 9.6 milling R1* 3.7 5.9 9.3 R2 5.1 9.2 16.5 R3* 3.6 6.6 12.0 R4 4.3 7.7 13.9 R5 3.7 7.0 15.5 R6* 4.3 8.1 16.5 R7 4.3 9.7 20.4 R8 5.3 10.8 21.3 R9* 3.1 5.5 9.3 R10 2.7 5.1 9.3 R11 3.3 5.5 8.7 R12 4.2 7.8 18.4 R13 4.2 6.8 10.8 R14 1.2 4.5 10.1

FIG. 7 shows particle size distributions (measured particle size being secondary particle size, not crystal or crystallite size) for samples 1, 2, 15, and 16, as a representative example of particle size distributions obtained by solid state routes in this study without further processing or size optimisation. The particle size distributions are typically bi-modal, with a first mode −10 μm, and a second mode −90 μm. Sample 3 presents significant differences in terms of particle size distribution, as shown in FIG. 8 due to the spray-drying and pyrolysis post-processing step.

All particle size distributions can also be refined with further processing steps, for example spray drying, ball milling, high shear milling, jet milling or impact milling to reduce the particle size distribution to the desired range (e.g. d90<20 μm, <10 μm or <5 μm) as shown in FIG. 15 and Table 2. Typically the particle size distributions are tuned by optimising the phase formation process (i.e. solid state synthesis route) and post-processing steps for the target application. For example, for a Li ion electrode with high power, one would typically target lower average particle sizes, amongst other considerations.

FIG. R5 shows the particle size distributions for samples R2, R4, R7, R10 in their final form, which are then processed into electrode slurries and inks.

SEM Characterisation of Samples

The morphology of some samples was analysed by Scanning Electron Microscopy (SEM).

FIGS. 9 and 10 show SEM images of sample 3 before and after pyrolysis. A porous microsphere morphology with carbon coating is observed, with primary crystallites organised into secondary particles.

It can be seen that the material has with homogeneous porous particles that can pack efficiently to form a high-density electrode. Qualitatively the conductivity is vastly improved as a conductive coating does not need to be applied for SEM imaging to be carried out, implying an order of magnitude improvement in material surface conductivity. FIG. 16 is an SEM image of the surface of a particle in an electrode of sample 22, where conductive carbon black particles contained in the electrode can also be seen in the right side of the image. This visibly shows evidence of a conformal carbon coating around the MNO material.

FIG. 11 shows SEM images of samples 1 and 2, and corroborates XRD and PSD data, showing compact secondary particle micron-size particles composed of ˜200 nm primary crystallites.

Electrochemical Testing of Samples

Electrochemical tests were carried out in half-coin cells (CR2032 size) for initial analysis. In half-coin tests, the material is tested in an electrode versus a Li metal electrode to assess its fundamental performance. In the below examples, the active material composition to be tested was combined with N-Methyl Pyrrolidone (NMP), carbon black acting as a conductive additive, and poly(vinyldifluoride) (PVDF) binder and mixed to form a slurry using a lab-scale centrifugal planetary mixer (although it is also possible to form aqueous slurries by using water rather than NMP, with binders such as CMC:SBR or alginate). The non-NMP composition of the slurries was 80 w. % active material, 10 w. % conductive additive, 10 w. % binder. The slurry was then coated on an Al foil current collector to the desired loading of 1 mg/cm² by doctor blade coating and dried in a vacuum oven for 12 hours. In this way, extremely thin coatings were achieved that enabled assessment of material fundamental properties, rather than those driven by electrode quality such as excess impedance, or poor packing of materials. Electrodes were punched out at the desired size and combined with a separator (Celgard porous PP/PE), Li metal, and electrolyte (1 M LiPF₆ in EC/DEC) inside a steel coin cell casing and sealed under pressure. Formation cycling was then carried out at low current rates (C/20) for 2 full charge and discharge cycles. After formation, further cycling can be carried out at a fixed or varied current density as required. These tests have been termed “half-cell galvanostatic cycling” for future reference. For samples R1-R10, the electrolyte was altered to 1.3 M LiPF₆ in 3:7 EC/DEC, and the formation cycling was carried out at C/10 for 2 charge/discharge cycles in the limits 1.1-3.0 V. The values shown for these samples is an average of 3 measurements, with the error being the standard deviation.

Homogeneous, smooth coatings on current collector foil, the coatings being free of visible defects were also prepared as above with a centrifugal planetary mixer to a composition of 94 w. % active material, 4 w. % conductive additive, 2 w. % binder. The coatings were calendared at 80° C. to a density of up to 3.0 g/cm³ at loadings of 1.3-1.7 mAh/cm² in order to demonstrate possible volumetric capacities >700 mAh/cm³ in the voltage range 0.7-3.0 V at C/20, and >640 mAh/cm³ in the voltage range 1.1-3.0 V at C/5. This is an important demonstration of these materials being viable in a commercially focussed electrode power cell formulation, where retaining performance after calendaring to a high electrode density allows for high volumetric capacities. Loadings of up to and including 1.0, 1.5, 2.0, 2.5, or 3.0 mAh/cm² may be useful for Li-ion cells focussed on power performance; loadings greater than 3.0 mAh/cm² are useful for energy-focussed performance in Li ion cells.

Electrical conductivity of electrodes made with the samples listed in Table 1 was measured using a 4-point probe thin film resistance measurement apparatus. Slurries were formulated according to the procedure described above and coated on a dielectric mylar film at a loading of 1 mg/cm². Electrode-sized discs where then punched out and resistance of the coated-film was measured using a 4-point probe. Bulk resistivity can be calculated from measured resistance using the following equation:

$\begin{matrix} {{{{{Bulk}{resistivity}(\rho)} = {2{{\pi s}\left( {V/I} \right)}}};}{{R = {V/I}};{s = {{0.1{cm}} = {2\pi \times 0.1 \times {R(\Omega)}}}}}} & (3) \end{matrix}$

The results of this test are shown in Table 3, below:

TABLE 3 Summary of 4-point probe resistivity measurement results for samples 1, 2, 4, 5, 6, 7, 13 to 20, and 22. Sample Resistance [kΩ] Bulk resistivity [kΩ · cm]  1* 8.5 5.3  2 1.7 1.1  4 3.2 2.0  5 0.52 0.33  6* 0.37 0.23  7 0.52 0.33 13 0.45 0.28 14 2.7 1.7 15 1.2 0.75 16 1.3 0.82 17 0.34 0.21 18 0.89 0.56 19 0.18 0.11 20 0.20 0.13 22 0.33 0.21

Samples R1-R14 also had their 4-point probe resistance measured to quantify their electrical resistivity. This was carried out with a different Ossila instrument (T2001A3-UK) at 2300 for coatings on mylar films at loadings of 1.0 mg/cm². The results for sheet resistance (Ω/square) are outlined in Table 3a, with error based on the standard deviation of 3 measurements.

TABLE 3a Summary of 4-point probe resistivity measurement results for samples R1 to R14. Sample Sheet Resistivity [Ω/square] R1*  1242 ± 156 R2  1041 ± 103 R3* 1396 ± 74 R4 1215 ± 52 R5 1057 ± 35 R6* 1092 ± 52 R7 1009 ± 89 R8  965 ± 83 R9* 1135 ± 92 R10 1113 ± 99 R12  891 ± 61 R13 1027 ± 13 12*  853 ± 51 R14  846 ± 57  6*  880 ± 29

The direct current internal resistance (DCIR) and the resultant area specific impedance (ASI) is a key measurement of internal resistance in the electrode in a Li-ion cell. In a typical measurement, a cell that has already undergone formation will be cycled at C/2 for 3 cycles. With the electrode in its delithiated state a C/2 discharge current is applied for 1 h to achieve ˜50% lithiation. The cell is rested for 30 mins to equilibrate at its OCV (open circuit voltage), and then a 5 C current pulse is applied for 10 s, followed by a 30 mins rest to reach the OCV. During the 10 s pulse the voltage response is sampled at a higher frequency to determine the average internal resistance accurately. The resistance is then calculated from V=IR, using the difference between the OCV (the linear average between the initial OCV before the pulse and afterwards) and the measured voltage. The resistance is then multiplied by the area of the electrode to result in the ASI.

The results of this test are shown in Table 4, below:

TABLE 4 Summary of DCIR/ASI measurement results for samples 1, 2, 4, 7, 14, 16, and 17. Sample ASI/Ω · cm²  1* 141  2 125  4 120  6* 126  7 162 13 67 14 99 16 74 17 162 18 75 19 164 22 121

The reversible specific capacity C/20, initial coulombic efficiency, nominal lithiation voltage vs Li/Li⁺ at C/20, 5C/0.5 C capacity retention, and 10 C/0.5 C capacity retention for a number of samples were also tested, the results being set out in Table 5, below. Nominal lithiation voltage vs Li/Li+ has been calculated from the integral of the V/Q curve divided by the total capacity on the 2^(nd) cycle C/20 lithiation. Capacity retention at 10 C and 50 has been calculated by taking the specific capacity at 100 or 5 C, and dividing it by the specific capacity at 0.5 C. It should be noted that the capacity retention was tested with symmetric cycling tests, with equivalent C-rate on lithiation and de-lithiation. Upon testing with an asymmetric cycling program, 10 C/0.5 C capacity retention greater than 89% is routinely observed.

Samples R1-R10 were tested with minor differences in Table 5a, the reversible specific capacity shown is the 2^(nd) cycle delithiation capacity at C/10, the nominal lithiation voltage vs Li/Li⁺ is at C/10 in the 2^(nd) cycle, the rate tests were carried out with an asymmetric cycling program with no constant voltage steps (i.e. constant current), with lithiation at C/5 and delithiation at increasing C-rates.

TABLE 5 Summary of electrochemical testing results from Li-ion half coin cells using a number of samples. In general (although not exclusively) it is beneficial to have a higher capacity, a higher ICE, a lower nominal voltage, and higher capacity retentions. Reversible specific Initial Nominal 5 C/0.5 C 10 C/0.5 C capacity coulombic lithiation capacity capacity C/20 efficiency voltage vs retention retention Sample [mAh/g] [%] Li/Li⁺ [V] [%] [%]  1* 214 87.8 1.61 62 35  2 240 90.9 1.61 64 45  3 203 84.9 1.58 79 68  4 286 90.7 1.59 68 54  5 253 86.0 1.60 63 43  6* 224 93.5 1.57 61 38  7 263 93.6 1.58 74 67  8* 192 82.0 1.60 54 36  9 188 86.8 1.61 64 54 10* 172 74.3 1.55 64 54 11 176 71.6 1.59 56 45 12* 164 93.9 1.77 86 81 13 184 95.4 1.75 86 80 14 278 91.0 1.59 15 228 89.2 1.59 16 281 90.8 1.58 72 58 17 203 94.6 1.58 18 228 90.1 1.59 84 68 19 193 87.0 1.56 63 44 21 169 70.9 1.59 67 56 22 267 86.9 1.57 71 62

TABLE 5a Summary of electrochemical testing results from Li-ion half coin cells using a number of samples. Specific capacity Initial Nominal ASI 5 C/0.5 C 10 C/0.5 C C/10 coulombic lithiation voltage [Ω · capacity capacity Sample [mAh/g] efficiency [%] C/10 [V] cm²] retention [%] retention [%] R1* 222 ± 7 98.23 ± 0.51 1.543 ± 0.001 169 ± 10 96.5 ± 0.1 95.9 ± 0.1 R2 273 ± 17 98.52 ± 0.45 1.550 ± 0.001 106 ± 18 97.3 ± 0.4 96.2 ± 0.7 R3* 244 ± 26 96.75 ± 0.31 1.549 ± 0.002 166 ± 17 96.1 ± 0.6 95.2 ± 0.8 R4 252 ± 9 98.80 ± 0.86 1.549 ± 0.001 109 ± 9 98.4 ± 0.0 97.4 ± 0.1 R5 272 ± 21 99.69 ± 1.56 1.549 ± 0.001 122 ± 3 96.3 ± 0.3 94.8 ± 0.4 R6* 134 ± 14 80.97 ± 1.55 1.539 ± 0.007 485 ± 75 72.8 ± 5.7 64.1 ± 7.2 R7 150 ± 8 82.15 ± 0.12 1.531 ± 0.000 390 ± 32 67.0 ± 0.4 56.8 ± 0.5 R8 144 ± 2 81.64 ± 1.35 1.530 ± 0.001 400 ± 42 72.9 ± 1.2 63.3 ± 1.5 R9* 211 ± 5 94.53 ± 0.18 1.630 ± 0.001 129 ± 13 96.2 ± 0.4 95.1 ± 0.5 R10 201 ± 7 98.42 ± 1.12 1.626 ± 0.000 118 ± 16 96.2 ± 0.1 94.9 ± 0.2 R12 198 ± 13 97.71 ± 0.25 1.544 ± 0.001 208 ± 8 95.2 ± 0.8 92.9 ± 1.0 R13 203 ± 15 98.22 ± 0.12 1.546 ± 0.001 199 ± 10 97.7 ± 0.0 97.7 ± 0.5

The modification of mixed niobium oxide-based Wadsley-Roth and Bronze structures as shown in the reference examples demonstrates the applicability of the modification to improve active material performance in Li-ion cells. By substituting the non-Nb cation to form a mixed cation structure as described, the entropy (cf disorder) can increase in the crystal structure, reducing potential energy barriers to Li ion diffusion through minor defect introduction (e.g. samples R7, 16). Modification by creating mixed cation structures that retain the same overall oxidation state demonstrate the potential improvements by altering ionic radii, for example replacement of an Mo⁶⁺ cation with W⁶⁺ in sample 14 or Fe³⁺ or Ga³⁺ for Al³⁺ in samples R4 and R5, which can cause minor changes in crystal parameters and Li-ion cavities (e.g. tuning the reversibility of Type VI cavities in Wadsley-Roth structures) that can improve specific capacity, Li-ion diffusion, and increase Coulombic efficiencies of cycling by reducing Li ion trapping. Modification by creating mixed cation structures that result in increased oxidation state (e.g. Ge⁴⁺ to replace Zn²⁺ in sample R2, or Mo⁶⁺ for Zr⁴⁺ in sample 19) demonstrate similar potential advantages with altered ionic radii relating to capacity and efficiency, compounded by introduction of additional electron holes in the structure to aid in electrical conductivity. Modification by creating mixed cation structures that result in decreased oxidation state (e.g. K⁺ and Co³⁺ to replace Ge⁴⁺ in sample R7, or Ti⁴⁺ to replace Mo⁶⁺ in sample 2) demonstrate similar potential advantages with altered ionic radii relating to capacity and efficiency, compounded by introduction of oxygen vacancies and additional electrons in the structure to aid in electrical conductivity. Modification by inducing oxygen deficiency from high temperature treatment in inert or reducing conditions demonstrate the loss of a small proportion of oxygen from the structure, providing a reduced structure of much improved electrical conductivity (e.g. sample 5, R10 and R12-14) and improved electrochemical properties such as capacity retention at high C-rates (e.g. sample 5, R13). Combination of mixed cation structures and induced oxygen deficiency allows multiple beneficial effects (e.g. increased specific capacity, reduced electrical resistance) to be compounded (e.g. samples 18, R8).

FIGS. 12, 13, and 17 show representative lithiation/delithiation curves for unmodified and modified MoNb₁₂O₃₃ (FIG. 12 —samples 1 and 6) ZrNb₂₄O₆₂ (FIG. 13 —samples 6 and 7), and W₇Nb₄O₃₁ (FIG. 17 —samples 12 and 13) in their first two formation cycles at C/20 rate. In FIG. 12 , approximately 90% of the specific capacity for sample 16 demonstrated is shown to be in a narrow voltage range of ca. 1.2-2.0 V, and in FIG. 13 approximately 90% of the capacity for sample 7 demonstrated is shown to be in a narrow range of ca. 1.25-1.75 V; these data highlight the attractive voltage profiles achievable with MNO crystals based upon Wadsley-Roth crystal structures. In FIG. 17 , approximately 90% of the specific capacity for sample 13 is shown to be in a narrow range of ca. 1.2-2.2 V; this demonstrates that attractive voltage profiles are achieved with MNO crystals based upon a tetragonal bronze crystal structure. Secondly, the complex metal oxide samples 7, 16, and 13 demonstrate improved specific capacity as compared to their unmodified crystals samples 1, 6 and 12. This is due to the cations that are included in the complex structures increasing the number of sites in the crystal that Li ions can accommodate due to their differing ionic radii and oxidation states, thus increasing capacity. An increase in ICE was observed between samples 1 and 16, and samples 12 and 13, which further demonstrates that Li ions intercalated in the modified crystal structure can be more efficiently delithiated as the Li ion sites are modified to enable their de-intercalation.

FIG. R5 demonstrates the particle size distribution of samples R2, R4, R8, R11 containing primarily a single peak that has a narrow distribution, i.e. D₁₀ and D₉₀ are similar in value to D₅₀. This is advantageous for processing the material in electrode slurries for efficient packing of the material, and to maintain a homogeneous electrochemical performance (e.g. a smaller particle will be fully lithiated in advance of a larger particle due to shorter diffusion distances).

FIG. R6 shows the advantage in modifying sample R1, particularly with regard to improving the observed specific capacity through substituting Zn²⁺ cations with Ge⁴⁺ cations of higher valency. FIG. R7 demonstrates the improved specific capacity observed on modifying sample R6 by substituting Ge⁴⁺ with K and Co cations, i.e. with cations of reduced valency. FIG. R9 demonstrates the improvement in ICE, and reduction in nominal lithiation voltage possible by introduction of induced oxygen vacancies that reduces polarisation effects through improving conductivity, and through improving the reversibility of lithiation/delithiation processes.

Across all materials tested, each modified (cation substituted and/or oxygen deficient) material demonstrates an improvement versus the unmodified ‘base’ crystal structure. This is inferred from measurements of resistivity/impedance by two different methods, and also electrochemical tests carried out in Li-ion half coin cells, particularly the capacity retention at increased current densities (cf. rates, Table 5). Without wishing to be bound by theory, the inventors suggest that this is a result of increased ionic and electronic conductivity of the materials as defects are introduced, or by alterations to the crystal lattice by varying ionic radii; also evidenced by DCIR/ASI (Table 4) and EIS (FIG. 14 ) measurements to show decreased resistance or impedance upon material modification. Li-ion diffusion rates likely also increase in modified materials, as compared with the unmodified ‘base’ materials. Specific capacities themselves may also increase in some cases as shown in Table 5, as doping/exchange with metal ions of different sizes can expand or contract the crystal lattice and allow for more intercalation or more reversibility of intercalation of Li-ions than possible in the unmodified structure.

The data in Table 3 show a large reduction in the resistivity between sample 1 (comparative) and samples 2, 4, 5, 14, 15, 16, 18, 22, demonstrating the effect of improving electrical conductivity of the crystal structures through both cation exchange, oxygen deficiencies, and carbon coating. Samples 17, 19, and 20 also show a similarly low resistivity versus sample 6. The resistivity slightly increased upon incorporation of 0.05 equivalents of V species in the base crystal in sample 7, however an improvement in specific capacity was observed due to the changes in available Li-ion sites in the crystal lattice likely as a result of the differing ionic radius of V over Zr (see Table 5).

The data in Table 4 shows a large reduction in the DCIR/ASI from sample 1 (comparative) to samples 2, 4, 14, 16, 18 and 22, reflecting the trends shown in Table 3. Samples 7, 17, and 19 demonstrate a higher than these by DCIR, however these relate to a different base crystal structure. Without wishing to be bound by theory, the inventors hypothesise that samples 7, 17, and 19 demonstrate an increase in DCIR/ASI as compared with the comparative material of sample 6 (ZrNb₂₄O₆₂) due to the changes in the crystal lattice with the introduced cations of different ionic radii. However, it remains beneficial in terms of conductivity for these structures for samples 17 and 19 as the electrical resistivity is decreased as shown in Table 3, thereby minimising joule heating and enabling a more uniform current distribution across the material, which in turn can enable improved safety and lifetime of a Li ion system. For sample 7, whilst there is no demonstrated improvement utilising V to exchange with Zr, there is an increase in specific capacity, as discussed above.

In Table 5, across most samples there is a trend for improved specific capacities, initial Coulombic efficiencies (ICE), nominal lithiation voltage vs Li/Li⁺, and importantly capacity retention at 5 C and 10 C vs 0.5 C for modified materials versus the comparative ‘base’ materials (e.g. samples 1, 6, 8, 10, 12). For example samples 2, 3, 4, 5, 14, 15, 16, 18, 22 all demonstrate improvements in one or more of these parameters vs sample 1. This is also the case for samples 7, 17, 19 versus sample 6 across multiple parameters; sample 11 and 21 versus 10 where an improvement in specific capacity or capacity retention is observed; sample 9 versus 8 where ICE and capacity retention are improved; and sample 13 versus 12 where ICE and capacity retention are improved.

Electrochemical impedance spectroscopy (EIS) measurements were also carried out to gain a further understanding on the impedance present in the electrode in a Li-ion cell. In a typical measurement, the cell is prepared as for DCIR measurements to ˜50% lithiation and then the frequency of alternating charge/discharge current pulses is varied whilst measuring the impedance. By plotting the real and imaginary components as the axes, and varying the AC frequency, a Nyquist plot is generated. From this plot for a Li-ion cell different types of impedance in the cell can be identified, however it is typically complex to interpret. For example, Ohmic resistance can be partially separated from electrochemical double layer effects and also separated from diffusion effects.

FIG. 14 (a) and (b) show EIS spectra for (comparative) sample 1 and samples 16 and 7 (modified samples).

EXAMPLES

Commercial-grade LTO (Li₄Ti₅O₁₂) was purchased from Targray Technology International Inc with properties outlined in Table E1 (Sample E1). The modified and carbon-coated Wadsley-Roth and Bronze materials were synthesised in-house by a solid-state route. In a first step precursor materials (e.g. Nb₂O₅, WO₃, ZrO₂, TiO₂, MoO₃, Cr₂O₃, ZnO, Al₂O₃ etc.) were mixed in stoichiometric proportions (200 g total) and ball-milled at 550 rpm with a ball to powder ratio of 10:1 for 3 h. The resulting powders were heat treated in an alumina crucible in a muffle furnace in air at T₁=800-1350° C. for 24 h, providing the desired Wadsley-Roth or Bronze phase. An additional heat treatment step was also applied under a N₂ atmosphere at T₂=800-1350° C. for 5 h to result in minor oxygen deficiencies in the base crystal structure for samples E2, E3, E4. For Sample E2 the above synthesis was carried out with T₁=9000, T₂=9000. For Sample E4 the above synthesis was carried out with T₁=11000, T₂=11000. For sample E5 the above synthesis was carried out with T₁=11000, repeated twice with an intermediary grinding step by impact milling at 20,000 rpm. For sample E6, the above synthesis was carried out with T₁=11000 for 24 h.

Sample E2 (98 g) was then combined with petroleum pitch (2 g) (ZL 118M available from Rain Carbon) by high energy impact mixing/milling. The mixture was heat treated in a furnace under reducing conditions at T=9000 for 5 h to provide Sample E3, which was a free-flowing black powder. A final de-agglomeration step was utilised for each sample by impact milling or jet milling to adjust to the desired particle size distribution. Specifically, the material was de-agglomerated by impact milling at 20,000 RPM for 10 seconds.

Active electrode material mixtures of MNO and LTO were obtained by low to high energy powder mixing/blending techniques, such as by rotational mixing in multiple directions, rotational V-type blending over a single axis, planetary mixing, centrifugal planetary mixing, high shear mixing, and other typical mixing/blending techniques. In this case, mixing was achieved with a centrifugal planetary mixer on 5 g batches of materials, mixed at 2000 rpm for 3 mins, 10 times.

TABLE E1 A summary of the materials utilised. Particle size distribution has been evaluated by dry powder laser diffraction, and surface area by the BET method using N₂. D10 D50 D90 BET Surface Sample Material (μm) (μm) (μm) Area [m² g⁻¹] E1 LTO (from commercial supplier) 0.8 2.5 5.0 16.0* E2 Ti_(0.05)Zr_(0.05)W_(0.25)Mo_(0.65)Nb₁₂O_(33−δ) 1.2 4.5 10.3 1.8 E3 Ti_(0.05)Zr_(0.05)W_(0.25)Mo_(0.65)Nb₁₂O_(33−δ) + C 1.8 4.3 8.6 1.6 E4 Ti_(0.35)W_(6.65)Nb₄O_(31−δ) 1.5 4.7 10.2 — E5 Al_(0.1)Zn_(1.9)Nb₃₄O_(87.05)** 4.0 6.1 9.2 — E6 Cr_(0.25)Al_(0.75)Nb₁₁O₂₉ 2.6 5.3 8.5 — *From manufacturer specification sheet. **Oxygen stoichiometry calculated assuming Al³⁺, Zn²⁺, Nb⁵⁺.

Materials Characterisation

The phase purity of samples was analysed using a Rigaku Miniflex powder X-ray diffractometer in 20 range (20-70°) at 1°/min scan rate.

FIG. 18 shows the measured XRD diffraction patterns for samples E1 through to E4. Diffractions patterns in Sample E1 has peaks at the same locations (within instrument error, that is 0.1°), and match JCPDS crystallography database entry JCPDS 49-0207, which corresponds to the spinel crystal structure of Li₄Ti₅O₁₂. There is no amorphous background noise and the peaks are sharp and intense. This means that the sample is crystalline, with crystallite size 43±7 nm according to the Scherrer equation. This confirms the presence of LTO with a spinel crystal structure.

Diffraction patterns in Sample E2 has peaks at the same locations (within instrument error, that is 0.10) and match JCPDS crystallography database entry JCPDS 73-1322, which corresponds to MoNb₁₂O₃₃. Sample E3 has some changes to its peaks due to the introduced oxygen-deficiency beginning to induce minor crystallographic distortions due to significant quantities of vacancy defects, and additional peaks relating to the carbon at ˜26° and ˜40°. There is no amorphous background noise and the peaks are sharp and intense. This means that all samples are crystalline, with crystallite size 38±4 nm for Sample E2 and 32±12 nm for Sample E3 according to the Scherrer equation and crystal structure matching MoNb₁₂O₃₃. This confirms the presence of a Wadsley-Roth crystal structure.

Diffraction patterns in Sample E4 has peaks at the same locations (within instrument error, that is 0.1°) and match JCPDS crystallography database entry database JCPDS 00-020-1320, which corresponds to W₇Nb₄O₃₁. There is no amorphous background noise and the peaks are sharp and intense. This means that the sample is phase-pure and crystalline, with crystallite size 43±10 nm according to the Scherrer equation and crystal structure matching W₇Nb₄O₃₁. This confirms the presence of a Tetragonal Tungsten Bronze crystal structure.

Sample E5 presented a phase mixture between the orthorhombic and monoclinic forms of the 3×4×∞ Wadsley-Roth structure, corresponding to crystallography database entry JCPDS 28-1478 and PDF card: 04-021-7859. There is no amorphous background noise and the peaks are sharp and intense. This means that the samples are phase-pure and crystalline, with crystallite size 49±6 nm according to the Scherrer equation and crystal structure matching Zn₂Nb₃₄O₈₇. This confirms the presence of a Wadsley-Roth crystal structure.

Diffraction patterns in Sample E6 have peaks at the same locations (with some shift due to crystal modification, up to around 0.2°), and match crystallography database entry JCPDS 22-009. There is no amorphous background noise and the peaks are sharp and intense. This means that the sample is crystalline, with crystallite size 45 nm according to the Scherrer equation and crystal structure matching AlNb₁₁O₂₉. This confirms the presence of a Wadsley-Roth crystal structure.

Thermogravimetric Analysis (TGA) was performed on Sample E3 using a Perkin Elmer Pyris 1 system in an air atmosphere. Samples were heated from 30° C. to 900° C. at 5° C./min, with an air flow of 20 mL/min.

TABLE E2 A summary of mass gain and mass loss as measured by TGA analysis in air. Sample Measured mass gain [%] Measured mass loss [%] E3 1.02 0.62

Particle Size Distributions were obtained with a Horiba laser diffraction particle analyser for dry powder. Air pressure was kept at 0.3 MPa. The results are set out in Table E1. BET surface area analysis was carried out with N₂ on a BELSORP-miniX instrument at 77.35 K and are set out in Table E1.

Electrochemical Characterisation

The electrochemical characterisation of the Examples was performed under different conditions to the Reference Examples. Therefore, in some instances the electrochemical characterisation of the Examples may not be directly comparable to the electrochemical characterisation of the Reference Examples.

Li-ion cell charge rate is usually expressed as a “C-rate”. A 1 C charge rate means a charge current such that the cell is fully charged in 1 h, 10 C charge means that the battery is fully charged in 1/10th of an hour (6 minutes). C-rate hereon is defined from the reversible capacity of the anode within the voltage limits applied, i.e. for an anode that exhibits 1.0 mAh cm⁻² capacity within the voltage limits of 1.1-3.0 V, a 1 C rate corresponds to a current density applied of 1.0 mA cm-2.

Electrochemical tests were carried out in half-coin cells (CR2032 size) for analysis. In half-coin tests, the active material is tested in an electrode versus a Li metal electrode to assess its fundamental performance. In the below examples, the active material composition to be tested was combined with N-Methyl Pyrrolidone (NMP), carbon black acting as a conductive additive, and poly(vinyldifluoride) (PVDF) binder and mixed to form a slurry using a lab-scale centrifugal planetary mixer. The non-NMP composition of the slurries was 90 wt % active material, 6 wt % conductive additive, 4 wt % binder. The slurry was coated on an Al foil current collector to the desired loading of 5.7-6.6 mg cm⁻² by doctor blade coating and dried. The electrodes were then calendared to a density of 2.00-3.75 g cm⁻³ (dependent on material density) at 80° C. to achieve targeted porosity of 35-42%. Porosity was calculated as the measured electrode density divided by the weighted average density of each component of the composite electrode coating film. Electrodes were punched out at the desired size and combined with a separator (Celgard porous PP/PE), Li metal, and electrolyte (1.3 M LiPF₆ in EC/DEC) inside a steel coin cell casing and sealed under pressure. Cycling was then carried out at low current rates (C/10) for 2 full cycles of lithiation and de-lithiation between 1.1-3.0 V. Afterwards, the cells were tested for their performance at increasing current densities. During rate tests, the cells were cycled asymmetric, with a slow lithiation (C/5, with a CV step at 1.1V to C/20 current density) followed by increasing de-lithiation rates for de-lithation rate tests. All electrochemical tests were carried out in a thermally controlled environment at 23° C.

The first cycle efficiency was calculated as the fraction of de-lithiation capacity/lithiation capacity in the first cycle at C/10. The nominal voltage at each C-rate was determined by integrating the voltage-capacity curves and then by dividing it by the total capacity.

To quantify the significance of the differences in data observed, an error calculation was carried out and applied to the values for specific capacity. The error for these was approximated as the largest error possible with the microbalance used (±0.1 mg), and the lowest loading electrode (5.7 mg cm⁻²) on a 14 mm electrode disc. This provides an error of t 1.1%, which has been applied to each capacity measurement. Error in Coulombic efficiency, capacity retention, and voltage were assumed to be negligible as the instrument accuracy far exceeds the stated significant figures, and the values are independent of the balance errors.

TABLE E3 A summary of the electrochemical tests undertaken with different mixtures of Samples E1 to E4. Achieved electrode conditions are also referenced for each test, providing smooth electrodes free of agglomerates, that demonstrate good adhesion and cohesion to the current collector. Test Ref. A* B* C* D* E F G H I J Content of Sample E1 [w/w %] 100 — — — 90 10 95 5 5 95 Content of Sample E2 [w/w %] — 100 — — 10 90 5 95 — — Content of Sample E3 [w/w %] — — 100 — — — — — — 5 Content of Sample E4 [w/w %] — — — 100 — — — — 95 — Electrode loading [mg cm⁻²] 6.3 5.7 6.6 6.4 6.5 6.1 6.1 6.4 6.2 6.3 *Comparative test of individual active electrode material

TABLE E4b A summary of the electrochemical tests undertaken with different mixtures of Samples E1, E5, E6. Achieved electrode conditions are also referenced for each test, providing smooth electrodes free of agglomerates, that demonstrate good adhesion and cohesion to the current collector. Test Ref. K* L* M N Content of Sample E1 [w/w %] — — 50 50 Content of Sample E5 [w/w %] 100 — 50 — Content of Sample E6 [w/w %] — 100 — 50 Electrode loading [mg cm⁻²] 6.4 6.4 6.3 6.6 *Comparative test of individual active electrode material

TABLE E5 A summary of electrochemical testing results from Li-ion half coin cells. In general (although not exclusively) it is beneficial to have a higher capacity, a higher ICE, and a lower area specific impedance. De-lithiation specific Initial coulombic Test capacity C/10 [mAh/g] efficiency [%] A* 161 ± 2 96.86 B* 211 ± 2 90.64 C* 221 ± 3 95.51 D* 164 ± 2 98.84 E 169 ± 2 91.25 F 201 ± 2 95.77 G 164 ± 2 96.13 H 213 ± 2 90.95 I 165 ± 2 98.15 J 164 ± 2 96.45 K* 207 ± 2 98.62 L* 202 ± 2 98.45 M 182 ± 2 97.69 N 185 ± 2 98.03 *Comparative test of individual active electrode material

TABLE E6 A summary of electrochemical testing results at increasing current densities from Li-ion half coin cells. In general (although not exclusively) it is beneficial to have a higher capacity retention. As these are measured in half-coin cells, the lithiation ability is severely limited at high C-rates due to limitations on Li ion extraction from the Li metal counter electrode, and so these results focus on de-lithiation ability. 1 C/0.5 C de- 2 C/0.5 C de- 5 C/0.5 C de- 10 C/0.5 C de- lithiation lithiation lithiation lithiation capacity capacity capacity capacity Test retention [%] retention [%] retention [%] retention [%] A* 99.4 98.8 97.5 96.3 B* 98.5 96.6 93.7 88.8 C* 98.6 96.8 95.0 91.0 D* 98.8 97.6 94.5 88.4 E 99.9 99.3 97.9 91.5 F 98.1 96.3 94.2 91.0 G 99.4 98.8 96.4 88.6 H 98.6 96.2 93.9 88.2 I 99.4 98.2 95.8 91.5 J 99.9 99.3 98.7 97.5 K* 99.6 99.0 98.1 87.8 L* 99.5 98.9 97.8 90.6 M 99.8 99.6 99.0 93.3 N 99.7 99.5 98.7 91.2 *Comparative test of individual active electrode material

TABLE E7 A summary of the nominal de-lithiation voltage at each C-rate. Nominal De-lithiation Voltage vs Li/Li⁺ [V] Test 0.1 C 0.5 C 1 C 2 C 5 C 10 C A* 1.57 1.59 1.60 1.61 1.67 1.78 B* 1.67 1.66 1.69 1.73 1.86 2.07 C* 1.61 1.61 1.63 1.67 1.78 1.94 D* 1.80 1.81 1.81 1.83 1.88 1.97 E 1.58 1.60 1.61 1.64 1.72 1.88 F 1.65 1.65 1.67 1.72 1.84 2.02 G 1.58 1.60 1.61 1.63 1.72 1.86 H 1.66 1.65 1.67 1.72 1.84 2.03 I 1.80 1.80 1.81 1.82 1.87 1.97 J 1.58 1.59 1.60 1.62 1.68 1.79 K* 1.59 1.61 — 1.70 1.85 2.09 L* 1.59 1.60 — 1.69 1.86 2.15 M 1.59 1.61 — 1.68 1.81 2.04 N 1.58 1.60 — 1.67 1.81 2.05 *Comparative test of individual active electrode material

Example A

Sample E2 has a Wadsley-Roth 3×4 block shear crystal structure based on a M^(VI)Nb₁₂O₃₃ crystal structure where all blocks are connected by tetrahedra, that has been made oxygen-deficient through heat treatment in an inert atmosphere and through cation exchange. The combination of induced oxygen-deficiency and cation exchange leads to improved electrochemical performance versus a material such as WNb₁₂O₃₃ or MoNb₁₂O₃₃.

As shown in tests A* and B*, Sample E2 has a higher specific capacity, lower ICE, lower capacity retention at higher C-rates, and higher nominal voltage at each C-rate than Sample E1. Therefore, by providing a physical mixture of the two materials in suitable proportions, the disadvantages of each can be alleviated. Due to the material design having suitable mixing characteristics, such as particle size distribution and surface chemistry, homogeneous powdered mixtures and subsequently homogeneous coated electrodes can be produced having an intimate mixture of the 2 components.

Tests E and G demonstrate mixtures with a high proportion of Sample E1 at 90 and 95 wt % respectively, and Tests F and H demonstrate mixtures with a high proportion of Sample E2 in a similar fashion. Tests E and G show increased specific capacity vs test A*, and improved initial Coulombic efficiency (ICE) vs test B; exemplifying the advantages of providing a physical mixture of LTO with Wadsley-Roth MNO materials. Tests F and H show increased specific capacity vs test A*, and improved ICE vs test B* in a similar fashion.

The de-lithiation capacity retention further demonstrates advantages to the mixture of active materials, with increased retention for tests E and G vs test B*. It is expected in a full cell arrangement with a cathode active material, that similar benefits will be observed for lithiation of the anode active material. The nominal de-lithiation voltage can be improved for the mixture of materials vs the individual active materials, with tests E, F, G, and H all showing a reduced nominal voltage compared to test B*.

FIGS. 19, 20, and 21 demonstrate some of the advantages of the mixtures discussed in their electrochemical characterisations in half-cells vs Li/Li⁺. FIG. 19 demonstrates the improved capacity for tests E and G over test A*, with the voltage curve shape being a combination of the individual tests A* and B. FIG. 20 demonstrates the reduced nominal voltage, and reduced observed polarisation of the composite electrode, for test F vs test B* at a high de-lithiation rate of 10 C. FIG. 21 demonstrates the improved voltage curve and reduced polarisation of test H vs test B at a rate of 5 C. These results importantly show that even small amount of mixed materials (e.g. test G, H with 5 wt %) can demonstrate significant advantages, with potential for even lower proportions of mixtures to allow precise tuning of electrochemical performance within desired parameters, such as specific capacity.

Sample E5 is a Wadsley-Roth 3×4 block shear crystal structure based on a M^(II) ₂Nb₃₄O₃₇ crystal structure composed of octahedra and no tetrahedra. Similar advantages can be observed in test M versus tests A* and K*.

Sample E6 is a Wadsley-Roth crystal structure based on M^(II)Nb₁₁O₂₉. Similar advantages can be observed in test N versus tests A* and L*. Notably, test N, a 50:50 mixture of Samples E1 and E6, was found to provide improved capacity retention at 1 C, 2 C, and 5 C compared to both tests A* and L*, the respective individual materials E1 and E6.

It is expected that similar benefits will be observed with all Wadsley-Roth crystal structures containing Nb as described in the claims mixed with lithium titanate as described above for use in Li-ion cells.

Example B

Sample E3 is a modified form of Sample E2, which has been coated with pitch-carbon by high energy milling, and then pyrolysed in an inert atmosphere to provide increased oxygen deficiency, and a polyaromatic sp²-based carbon coating based on a pitch precursor. This provides advantages in reducing impedance, reducing nominal voltage, and improving performance at high rate vs Sample E2. It carries further advantages such as improved surface electrical conductivity of the active material crystal, and improved mixing with other components of the electrode such as the carbon additive (typically carbon black, graphite, etc).

Test J shows an improved specific capacity vs test A*, and an improved ICE vs test C. The de-lithiation capacity retention is greater for test J than for both test A* and C*, implying that providing the mixture of these 2 different active materials is advantageous for high rate performance over either individual materials of Sample E1 and E3. This could be due to a combination of favourable surface chemistry of Sample E1 and E3 leading to enhanced electrode quality (adhesion, cohesion, conductivity), or a favourable combination of the electrochemical properties that can prevent impedance more effectively than the individual materials. Furthermore, the nominal de-lithiation voltage is decreased for test J vs test C*.

FIG. 23 demonstrates the improved capacity for test J over test A*, with the voltage curve shape being a combination of the individual tests A* and C*.

It is expected that similar benefits will be observed with all MNO materials that are coated with carbon and are oxygen-deficient as described mixed with lithium titanate for use in Li-ion cells.

Example C

Sample E4 has a Bronze crystal structure that has been made partially oxygen-deficient by a heat treatment in an inert atmosphere and by cation exchange. Specifically, the M^(VI) ₇Nb₄O₃₁ base crystal structure has 3, 4, and 5 sided tunnels with a low degree of filled tunnels, resulting in a high availability of Li-ion intercalation sites. In this case the combination of cation exchange and oxygen deficiency provides improved electrochemical performance versus materials such as W₇Nb₄O₃₁.

Test I demonstrates increased specific capacity and increased ICE versus test A*. Test I further demonstrates great improvement in de-lithiation capacity retention vs test D, more so than expected with inclusion of Sample E1 as the minor component (5 w/w %). FIG. 22 shows the advantage in capacity retention observed at a rate of 10 C for test I over test D*.

It is expected that similar benefits will be observed with all Bronze (Tetragonal Tungsten Bronze) crystal structures containing Nb as described in the claims mixed with lithium titanate for use in Li-ion cells.

While the invention has been described in conjunction with the exemplary embodiments described above, many equivalent modifications and variations will be apparent to those skilled in the art when given this disclosure. Accordingly, the exemplary embodiments of the invention set forth above are considered to be illustrative and not limiting. Various changes to the described embodiments may be made without departing from the spirit and scope of the invention.

For the avoidance of any doubt, any theoretical explanations provided herein are provided for the purposes of improving the understanding of a reader. The inventors do not wish to be bound by any of these theoretical explanations.

Any section headings used herein are for organizational purposes only and are not to be construed as limiting the subject matter described.

REFERENCES

A number of publications are cited above in order to more fully describe and disclose the invention and the state of the art to which the invention pertains. Full citations for these references are provided below.

The entirety of each of these references is incorporated herein.

-   [1] J. B. Goodenough et. al., J. Am. Chem. Soc., 135, (2013),     1167-1176. -   [2] R. J. Cava, J. Electrochem. Soc., (1983), 2345. -   [3] R. J. Cava, Solid State Ionics 9 & 10 (1983)407-412 -   [4] Kent J. Griffith et. al., J. Am. Chem. Soc., 138, (2016),     8888-8889. -   [5] Kent J. Griffith et. al., Inorganic Chemistry., 56, (2017),     4002-4010. -   [6] Sagrario M. Montemayor et. al., J. Mater. Chem., 8 (1998),     2777-2781. -   [7] Botella et. Al., Catalysis Today, 158 (2010), 162-169. 

1. An active electrode material comprising a mixture of (a) at least one lithium titanium oxide and (b) at least one mixed niobium oxide, wherein the mixed niobium oxide is expressed by the formula [M1]_(x)[M2]_((1-x))[Nb]_(y)[O]_(z), wherein: M1 and M2 are different; M1 is selected from one or more of P, B, Ti, Mg, V, Cr, W, Zr, Mo, Cu, Fe, Ga, Ge, Ca, K, Ni, Co, Al, Sn, Mn, Ce, Se, Si, Sb, Y, La, Hf, Ta, Zn, In, and Cd; M2 is selected from one or more of P, Mg, V, Cr, W, Zr, Mo, Cu, Fe, Ga, Ge, Ca, K, Ni, Co, Al, Sn, Mn, Ce, Sb, Bi, Sr, Y, La, Hf, Zn, Ta, In, and Cd; and wherein x satisfies 0≤x<0.5; y satisfies 0.5≤y≤49; z satisfies 4≤z≤124; with the proviso that if x=0 and M2 consists of a single element then the mixed niobium oxide is oxygen deficient.
 2. The active electrode material according to claim 1, wherein (i) M2 is selected from one or more of Mo, W, V, Zr, P, Al, Zn, Ga, Ge, Ta, Cr, Cu, K, Mg, Ni, and Hf; or (ii) M2 is selected from one or more of Mo, W, V, Zr, P, Al, Zn, Ga, and Ge; or (iii) M2 is selected from one or more of Mo, W, V, and Zr.
 3. The active electrode material according to any preceding claim, wherein M1 has an equal or lower oxidation state than M2, optionally wherein M1 has a lower oxidation state than M2.
 4. The active electrode material according to any preceding claim, wherein M1 comprises at least one cation with a 4+ oxidation state and wherein M2 comprises at least one cation with a 6+ oxidation state; optionally wherein M1 has an oxidation state of 4+ and wherein M2 has an oxidation state of 6+.
 5. The active electrode material according to any preceding claim, wherein (i) M1 is selected from one or more of P, B, Ti, Mg, V, Cr, W, Zr, Mo, Cu, Fe, Ga, Ge, K, Ni, Co, Al, Hf, Ta, and Zn; or (ii) M1 is selected from one or more of P, B, Ti, Mg, V, Cr, W, Zr, Mo, Ga, Ge, Al, and Zn; or (iii) M1 is selected from one or more of Ti, Zr, V, W, and Mo.
 6. The active electrode material according to any preceding claim, wherein (i) x satisfies 0<x<0.5; and/or (ii) x satisfies 0.01<x<0.4; and/or (iii) x satisfies 0.05≤x≤0.25.
 7. The active electrode material according to any preceding claim, wherein (i) the ratio by mass of (a):(b) ranges from 0.5:99.5 to 99.5:0.5; and/or (ii) the ratio by mass of (a):(b) ranges from 2:98 to 98:2; and/or (iii) the ratio by mass of (a):(b) is at least 2:1, at least 5:1, or at least 8:1; or (iv) the ratio by mass of (b):(a) is at least 2:1, at least 5:1, or at least 8:1.
 8. The active electrode material according to any preceding claim, wherein (i) the lithium titanium oxide has a spinel or ramsdellite crystal structure; and/or (ii) the lithium titanium oxide has a crystal structure as determined by X-ray diffraction corresponding to Li₄Ti₅O₁₂ and/or Li₂Ti₃O₇; and/or (iii) the lithium titanium oxide is selected from Li₄Ti₅O₁₂, Li₂Ti₃O₇, and mixtures thereof.
 9. The active electrode material according to any preceding claim, wherein (i) the lithium titanium oxide is doped with additional cations or anions; and/or (ii) the lithium titanium oxide is oxygen deficient; and/or (iii) the lithium titanium oxide comprises a coating, optionally wherein the coating is selected from carbon, polymers, metals, metal oxides, metalloids, phosphates, and fluorides.
 10. The active electrode material according to any preceding claim, wherein the crystal structure of the mixed niobium oxide as determined by X-ray diffraction corresponds to the crystal structure of the unmodified form of the mixed niobium oxide, wherein the unmodified form is expressed by the formula [M2][Nb]_(y)[O]_(z) wherein M2 consists of a single element and wherein the unmodified form is not oxygen deficient, wherein the unmodified form is selected from M2^(I)Nb₅O₁₃, M2^(I) ₆Nb_(10.8)O₃₀, M2^(II)Nb₂O₆, M2^(II) ₂Nb₃₄O₈₇, M2^(III)Nb₁₁O₂₉, M2^(III)Nb₄₉O₁₂₄, M2^(IV)Nb₂₄O₆₂, M2^(IV)Nb₂O₇, M2^(IV) ₂Nb₁₀O₂₉, M2^(IV) ₂Nb₁₄O₃₉, M2^(IV)Nb₁₄O₃₇, M2^(IV)Nb₆O₁₇, M2^(IV)Nb₁₈O₄₇, M2^(V)Nb₉O₂₅, M2^(V) ₄Nb₁₈O₅₅, M2^(V) ₃Nb₁₇O₅₀, M2^(VI)Nb₁₂O₃₃, M2^(VI) ₄Nb₂₆O₇₇, M2^(VI) ₃Nb₁₄O₄₄, M2^(VI) ₅Nb₁₆O₅₅, M2^(VI) ₈Nb₁₈O₆₉, M2^(VI)Nb₂O₈, M2^(VI) ₁₆Nb₁₈O₉₃, M2^(V) ₂₀Nb₂₂O₁₁₅, M2^(VI) ₉Nb₈O₄₇, M2^(VI) ₈₂Nb₅₄O₃₈₁, M2^(VI) ₃₁Nb₂₀O₁₄₃, M2^(VI) ₇Nb₄O₃₁, M2^(VI) ₁₅Nb₂O₅₀, M2^(VI) ₃Nb₂O₁₄, and M2^(VI) ₁₁Nb₁₂O₆₃, wherein the numerals I, II, III, IV, V, and VI represent the oxidation state of M2.
 11. The active electrode material according to any preceding claim, wherein the mixed niobium oxide is selected from: M1_(x)Mo_((1-x))Nb₁₂O_((33-33α)) M1_(x)W_((1-x))Nb₁₂O_((33-33α)) M1_(x)Mo_((1-x))Nb_(4.667)O_((14.667-14.667α)) (i.e. Mo₃Nb₁₄O₄₄ base structure) M1_(x)V_((1-x))Nb₉O_((25-25α)) M1_(x)Zr_((1-x))Nb₂₄O_((52-52α)) M1_(x)Zn_((1-x))Nb₁₇O_((43.5-43.5α)) (i.e. Zn₂Nb₃₄O₈₇ base structure) M1_(x)Cu_((1-x))Nb₁₇O_((43.5-43.5α)) (i.e. Cu₂Nb₃₄O₈₇ base structure) M1_(x)W_((1-x))Nb_(0.571)O_((4.429-4.429α)) (i.e. W₇Nb₄O₃₁ base structure) M1_(x)W_((1-x))Nb_(0.889)O_((5.222-5.222α)) (i.e. W₉Nb₈O₄₇ base structure) M1_(x)W_((1-x))Nb_(3.2)O_((11-11α)) (i.e. W₅Nb₁₅O₅₅ base structure) M1_(X)W_((1-x))Nb_(1.125)O_((5.813-5.813α)) (i.e. W₁₆Nb₁₈O₉₃ base structure) M1_(x)Al_((1-x))Nb₁₁O_((29-29α)) M1_(x)Ga_((1-x))Nb₁₁O_((29-29α)) M1_(x)Fe_((1-x))Nb₁₁O_((29-29α)) M1_(x)Al_((1-x))Nb₄₉O_((124-124α)) M1_(x)Ga_((1-x))Nb₄₉O_((124-124α)) M1_(x)Fe_((1-x))Nb₄₉O_((124-124α)) M1_(x)Ge_((1-x))Nb₁₃O_((47-47α)) wherein α satisfies 0≤α≤0.05 wherein x and/or α is >0.
 12. The active electrode material according to any preceding claim, wherein the mixed niobium oxide is oxygen deficient, optionally wherein z is defined as z=(z′−z′α) wherein α satisfies 0<α≤0.05.
 13. The active electrode material according to any preceding claim, wherein the mixed niobium oxide is oxygen deficient and is selected from: MoNb₁₂O_((33-33α)) WNb₁₂O_((33-33α)) MO₃Nb₁₄O_((44-44α)) VNb₉O_((25-25α)) ZrNb₂₄O_((62-62α)) Zn₂Nb₃₄O_((87-87α)) Cu₂Nb₃₄O_((87-87α)) W₇Nb₄O_((31-31α)) W₉Nb₈O_((47-47α)) W₅Nb₁₆O_((55-55α)) W₁₆Nb₁₈O_((93-93α)) AlNb₁₁O_((29-29α)) GaNb₁₁O_((29-29α)) FeNb₁₁O_((29-29α)) AlNb₄₉O_((124-124α)) GaNb₄₉O_((124-124α)) FeNb₄₉O_((124-124α)) GeNb₁₈O_((47-47α)) wherein α satisfies 0<α≤0.05.
 14. The active electrode material according to any preceding claim, wherein the mixed niobium oxide has a Wadsley-Roth crystal structure and/or a Tetragonal Tungsten Bronze crystal structure.
 15. The active electrode material according to any preceding claim, wherein the lithium titanium oxide is in particulate form, optionally wherein the lithium titanium oxide has a Doo particle diameter in the range of 0.1-50 μm, or 0.25-20 μm, or 0.5-15 μm.
 16. The active electrode material according to any preceding claim, wherein the mixed niobium oxide is in particulate form, optionally wherein the mixed niobium oxide has a D₅₀ particle diameter in the range of 0.1-100 μm, or 0.5-50 μm, or 1-25 μm.
 17. The active electrode material according to any preceding claim, wherein the lithium titanium oxide and the mixed niobium oxide are in particulate form and wherein the ratio of the D₅₀ particle diameter of the lithium titanium oxide to the D₅₀ particle diameter of the mixed niobium oxide is in the range of 0.01:1 to 0.9:1, or 0.1:1 to 0.7:1.
 18. The active electrode material according to any preceding claim, wherein the lithium titanium oxide has a BET surface area in the range of 0.1-100 m²/g, or 1-50 m²/g, or 3-30 m²/g.
 19. The active electrode material according to any preceding claim, wherein the mixed niobium oxide has a BET surface area in the range of 0.1-100 m²/g, or 0.5-50 m²/g, or 1-20 m²/g.
 20. The active electrode material according to any preceding claim, wherein the ratio of the BET surface area of the lithium titanium oxide to the BET surface area of the mixed niobium oxide is in the range of 1.1:1 to 20:1, or 1.5:1 to 10:1.
 21. The active electrode material according to any preceding claim, wherein the mixed niobium oxide is coated with carbon, optionally wherein (i) the coating comprises polyaromatic sp² carbon; and/or (ii) the coating is derived from pitch carbons; and/or (iii) wherein the coating is present in an amount of up to 10 wt %, or 0.05-5 wt %, or 0.1-3 wt %, based on the total weight of the mixed niobium oxide and the coating.
 22. The active electrode material according to any preceding claim, wherein the crystal structure of the mixed niobium oxide, as determined by X-ray diffraction analysis, corresponds to the crystal structure of one or more of: MoNb₁₂O₃₃ WNb₁₂O₃₃ Mo₃Nb₁₄O₄₄ VNb₉O₂₅ ZrNb₂₄O₆₂ Zn₂Nb₃₄O₇ Cu₂Nb₃₄O₃₇ W₇Nb₄O₃₁ W₉Nb₈O₄₇ W₅Nb₁₆O₅₅ W₁₅Nb₁₈O₉₃ AlNb₁₁O₂₉ GaNb₁₁O₂₉ FeNb₁₁O₂₉ AlNb₄₉O₁₂₄ GaNb₄₉O₁₂₄ FeNb₄₉O₁₂₄ GeNb₁₈O₄₇.
 23. The active electrode material according to any preceding claim, wherein the lithium titanium oxide and/or the mixed niobium oxide further comprises lithium and/or sodium.
 24. A composition comprising the active electrode material of any preceding claim and at least one other component; optionally wherein at least one other component is selected from a binder, a solvent, a conductive additive, an additional active electrode material, and mixtures thereof.
 25. An electrode comprising the active electrode material of any of claims 1-23 in electrical contact with a current collector.
 26. An electrochemical device comprising an anode, a cathode, and an electrolyte disposed between the anode and the cathode, wherein the anode comprises an active electrode material according to any of claims 1-23; optionally wherein the electrochemical device is a lithium-ion battery or a sodium-ion battery.
 27. A method for making an active electrode material, wherein the active electrode material is as defined in any of claims 1-23, the method comprising mixing at least one lithium titanium oxide with at least one mixed niobium oxide. 